All-Solid-State Thin Film Li-Ion Batteries: New Challenges, New Materials, and New Designs

Abstract

All-solid-state batteries (ASSBs) are among the remarkable next-generation energy storage technologies for a broad range of applications, including (implantable) medical devices, portable electronic devices, (hybrid) electric vehicles, and even large-scale grid storage. All-solid-state thin film Li-ion batteries (TFLIBs) with an extended cycle life, broad temperature operation range, and minimal self-discharge rate are superior to bulk-type ASSBs and have attracted considerable attention. Compared with conventional batteries, stacking dense thin films reduces the Li-ion diffusion length, thereby improving the rate capability. It is vital to develop TFLIBs with higher energy density and stability. However, multiple challenges, such as interfacial instability, low volumetric energy density, and high manufacturing cost, still hinder the widespread application of TFLIBs. At present, many approaches, such as materials optimization and novel architecture design, have been explored to enhance the stability and energy density of TFLIBs. An overview of these discoveries and developments in TFLIBs is presented in this review, together with new insights into the intrinsic mechanisms of operation; this is of great value to the batteries research community and facilitates further improvements in batteries in the near future.

1. Introduction

Lithium-ion batteries (LIBs) have become ubiquitous in our modern lives, with various applications, such as portable electronic devices, electric vehicles, and grid storage [1,2,3]. In conventional LIBs, liquid electrolytes are favored for their high electrochemical stability voltage window, high ionic conductivity, and excellent wettability with other internal composites of batteries [4]. Liquid electrolytes are typically based on a mixture of carbonate-based organic solvents with Li-salts and additives. Liquid electrolytes face significant drawbacks, such as high flammability, limited operating temperature range, and the possibility of irreversible decomposition [5]. In addition, forming Li-dendrites in organic liquid electrolytes will cause internal short circuits, resulting in the complete failure of the whole battery [6,7].

All-solid-state batteries (ASSBs), with non-volatile solid-state electrolytes (SSEs), present a viable alternative to replace traditional liquid-electrolyte LIBs [8], which can improve the energy density [9,10,11,12]. ASSBs are categorized into two types: bulk and thin film types. Compared to bulk-type ASSBs, all-solid-state thin film LIBs (TFLIBs) permit higher charge/discharge rates thanks to the improved contact between the thin film electrodes and SSEs [13].

TFLIBs have been used for many low-power microelectronic devices as onboard energy supply [14,15]. Compared with the porous electrodes in bulk-type ASSBs, the thin film electrode materials are very dense. In addition, as the electrolyte and electrode of the TFLIBs are deposited on an atomic/molecular level, various microscopic defects at the solid-solid interfaces can be avoided, which implies fewer interfacial issues. In addition to the inherent benefits of ASSBs, such as high energy density, a wide operating temperature range, and long cycle life, TFLIBs also have the following key features in comparison to bulk-type ASSBs [16]: (1) the electrode/electrolyte interfaces are in good contact, and the electrolyte layer is very thin, which facilitates fast charging and discharging; (2) the electrode material is more compact and can therefore achieve higher energy densities, a lower self-discharge rate (<1% per year), and a longer cycle life; (3) small-sized batteries can be designed to match semiconductor production processes, and can be integrated into electronic chips. Therefore, TFLIBs can be widely used in smart cards, electronic tags, integrated circuits, medical implants, and other fields. It is regarded as an enormously prospective energy source for microelectronics and a fantastic way of powering wearable electronic devices [17].

With the continuous advancements, TFLBs have become a promising power supplier for many devices owing to their compact size, flexibility, and impressive energy density [18,19,20,21]. Some of the most promising applications of TFLIBs include [22,23]: (1) Wearable Devices: TFLIBs are the ideal power source for wearable devices such as smartwatches, fitness trackers, and medical sensors due to their small size, flexibility, and lightweight; (2) Internet of Things Devices: TFLIBs can power internet of things devices that require small, low-power batteries, such as smart home sensors, wireless tags, and environmental monitoring devices; (3) Medical Implants: TFLIBs can power medical implants, such as pacemakers and defibrillators, in which a small size and long-term reliability are required; (4) Military and Aerospace Applications: TFLBs are suitable for military and aerospace applications owing to their lightweight construction, small size, and high energy density; (5) Energy Harvesting: TFLBs can be used in energy harvesting applications, where they can store energy generated by solar panels, thermoelectric generators, or other energy sources; (6) Smart Cards: TFLBs can be used in smart cards for secure transactions, such as credit cards, passports, and other identification documents. Understanding the correlation of the performance with the materials composition, solid-state interfaces, and battery architectures will provide insight into enhancing the performance of TFLIBs. However, TFLBs face significant challenges arising from the permanent growth in demands of electrical devices, such as ultra-long life, high capacity, and fast charging rates. However, more issues have to be resolved to reach: (1) High specific electrode storage capacities and ionic conductivities of SSEs; (2) Good interface stability and low interfacial resistance between inorganic SSEs and electrodes; (3) High voltage and good rate performance of TFLIBs.

Although serval reviews have recently been published to report the progress of TFLIBs [18,19,20,21], the majority of the reviews have been directed towards the recent progress in SSEs and electrode materials. Few reviews have discussed the advanced architecture design of TFLIBs to increase the energy density and reduce the manufacturing costs. However, these components are crucial for the practical applications of TFLIBs [22,23].

This review outlines the recent advancements and achievements of new constituent materials and architectural designs in TFLIBs. The historical developments and current challenges of TFLIBs are presented first. The following section summarizes the novel constituent materials developed to improve the TFLIBs’ performance, including electrode materials, electrolytes, and current collectors. Then, the novel designs of TFLIBs’ architectures, developed in recent years to achieve higher energy densities, are described. Finally, future research opportunities in constituent materials and manufacturing technologies are discussed. The prospect of novel architecture designs for high-energy TFLIBs, including advanced planar TFLIBs, three-dimensional (3D) TFLIBs, and other novel structural TFLIBs, are outlined. Such a thorough evaluation will help to guide future developments of TFLIBs with a high specific capacity, high safety, and extended cycle life.

2. History and Challenges

2.1. Brief History of TFLIBs

TFLIBs are deposited by thin film deposition methods, such as physical vapor deposition (PVD), vacuum evaporation, chemical vapor deposition (CVD), or atomic layer deposition (ALD) [24]. TFLIBs are typically manufactured on a solid substrate, such as glass, ceramic, or even polymer. The construction typically comprises a current collector layer, an anode layer, an electrolyte layer, a cathode layer, and another current collector layer. Sator reported the first thin film cell in 1952 [25]; it featured a lead chloride electrolyte deposited by vacuum evaporation. Then, the first Li-ion thin film batteries (AgI||LiI||Li) were reported in 1969 [26]. Over the next 20 years, the primary focus of research was on enhancing the performance of SSEs and electrode materials. As a result, TFLIBs were rapidly developed. In 1994, C/LiCoO₂ (LCO) batteries had a potential exceeding 3.6 V and boasted an impressive energy density, ranging between 120 and 150 Wh/kg (2 to 3 times that of ordinary NiCd batteries), which were successfully commercialized, and were considered the highest-performance portable power device [27]. Since then, LiCoO₂ has become the most popular and widely used cathode material. Another milestone was SSE films of lithium phosphorus oxynitride (LiPON), prepared by Bates in 1992 [28]. At present, LiPON is widely acknowledged as a conventional solid electrolyte utilized in TFLIBs. Subsequently, a series of microscale TFLIBs have been gradually designed. In 2002, the microscale TFLIBs (LiCoO₂||LiPON||Li) were reported to deliver a highly stable volumetric capacity of 40 µAh cm⁻²·µm⁻¹, operating at 3.9 V [29]. To further enhance the energy density of TFLIBs, the 3D integrated TFLIBs (LiCoO₂||LiPON||Si) were proposed in 2007 [30]. The utilization of a 3D structural design is promising for increasing the footprint capacity and volumetric capacity density. Then, various promising electrode materials and solid electrolytes were gradually explored for TFLIBs to improve their stability and capacity. In 2015, Cras et al. reported the microscale TFLIBs (Li_1.2TiO_0.5S_2.1||LiPON||Si), which demonstrated a cycled volumetric capacity of approximately 38 µAh·cm⁻²·µm⁻¹ over 740 cycles [31]. In 2017, STMicroelectronics proposed TFLIBs of LiCoO₂||LiPON||Li achieving a nominal voltage of 3.9 V and a capacity of 1 mAh [29]. In 2018, Rupp et al. used PLD to deposit the Li₇La₃Zr₂O₁₂ (LLZO) film and achieved promising solid electrolyte films for future TFLIBs by changing the deposition temperature to enhance the ionic conductivity [32]. In 2020, TDK Electronics introduced an all-ceramic structured TFLB with a capacity of 100 μAh and a nominal voltage of 1.5 V [18]. In addition to the mentioned developments, several companies, such as Infinite Power Solutions, Cymbet, and EnFilm, have initiated the commercialization of TFLBs for powering microelectronic devices.

2.2. Current Challenges

With the development of new manufacturing processes and the explosive growth in new chemistries, many promising materials have been investigated for TFLIBs (Figure 1). However, few TFLIBs have been commercialized. Mismatches in the voltage range, chemical reactions between various battery components, and high solid-state interface impedances result in the batteries possessing a lower capacity, cycle life, and energy density. In addition, as electronic devices continue to rapidly be developed, there is an increasing demand for batteries with improved performance, such as faster charging rates, larger capacities, higher energy densities, and better cycling performance [33]. This section reviews the main challenges and improvements for the current TFLIBs.

2.2.1. Materials and Deposition Technologies

The properties of the materials utilized significantly impact the performance of batteries. The SSE film plays a critical role in the safety and electrochemical performance of TFLIBs, which is a bottleneck for the development of TFLIBs. A high-performance SSE should show high ionic and low electronic conductivity, induce low electrode corrosion, and have a wide electrochemical operating window, etc. In order to obtain high ionic conductivity, the composition and structure of the electrolytes must be adjustable [20]. More importantly, the electrolyte must be compatible with the electrode materials, forming a stable solid/electrolyte interface. Li₃PO₄ and LIPON are still commonly used solid electrolytes, but their conductivity and stability are insufficient for high-energy batteries. Exploring solid electrolyte films with high ionic conductivity, high compatibility, broad electrochemical stability windows, and cycling stability is imminent [34].

In addition, electrode materials also face significant challenges in improving the batteries’ performance. Safety, theoretical capacity, cost, and recyclability must be carefully considered when selecting electrode materials. However, the current electrode materials significantly reduce the electrode materials’ sustainability because they contain rare elements and are produced by complex ceramic processing. Furthermore, the actual energy cost of battery manufacturing must include the overall cost of the electrode synthesis and manufacturing equipment [35,36].

Furthermore, negative electrodes should preferably have a low reduction potential. In contrast, positive electrodes should exhibit a high oxidation potential to ensure a high output voltage. Although many novel electrode materials have been synthesized, such as carbon and LiCoO₂, most of them excel in only a few aspects, which cannot really meet all of the requirements for high-performance TFLIBs [37]. For example, carbon and Li-metal are the most widely used anode materials. Carbon combines high stability and conductivity, but its low theoretical capacity limits its application. Li-metal anodes have a high storage capacity, but their poor stability results in low recyclability. Combining these performance requirements is crucial in exploring better electrode materials.

Thin film deposition technologies are also a great challenge for the fabrication of TFLIBs. The quality and the composition homogeneity of the films directly affect the battery’s performance. According to the deposition mechanism, the methods to fabricate TFLIBs can be roughly classified into CVD and PVD. PVD involves evaporating the material source in a vacuum to create gaseous atoms, molecules, or ions [38]. The thin film, with specific functions, is deposited on the substrate surface using a low-pressure gas or plasma. The most common PVD methods are vacuum evaporation, sputtering, and ion plating. PVD can deposit metals, alloys, compounds, ceramics, polymer films, and semiconductors. CVD is a method that uses one or more gaseous compounds or simple substances to form thin films on substrate surfaces by chemical reaction. CVD can be used to deposit a variety of inorganic materials, including single-crystal, polycrystalline, or glassy thin films.

CVD is typically carried out using different techniques, such as plasma-enhanced CVD (PECVD), metal-organic CVD (MOCVD), and low-pressure CVD (LPCVD), etc. [20]. However, as the component materials include increasingly constituent elements and complex structures, the battery’s design becomes more complex. The mentioned deposition techniques are insufficient to produce such films. In addition, many powder materials with better performance have been explored recently, such as lithium lanthanum titanates (LLT) [39], Li₇La₃Zr₂O₁₂ (LLZO) [40], LiNi_xCo_yMn_zO₂ [41], etc. However, due to their complex structure and composition, the quality and homogeneity of the corresponding films cannot be easily controlled by the above deposition methods. Therefore, innovation in deposition techniques is needed, along with exploring new active and conductive materials. Fortunately, the recent development of novel deposition technologies, such as ALD, PLD, and 3D printing, will provide more advanced methods for depositing complex thin film battery materials.

2.2.2. Solid-Solid Interfaces

Although the performance of the individual battery components (e.g., electrodes or electrolytes) has been improved considerably, the performance improvement of the whole cell is still lean. Moreover, the power density and lifetime of TFLIBs do not yet meet the practical application requirements. The low performance is mainly attributed to the significant interface resistance that exists between the electrodes and SSEs [42]. This includes the following aspects, depicted in Figure 2: poor contact, formation of an interlayer, and interfacial reactions.

The interface resistance between the electrodes and electrolyte largely depends on the change in the electrode volume during the cycling process [43]. This would lead to poor contact between the electrodes and electrolyte, which can eventually result in crack formation [44] and interfacial delamination [45,46]. Insufficient mechanical contacts can also lead to “dead corner” areas caused by electrode contacts separated by solid electrolytes [47]. The absence of conductive pathways inhibits the electrons and Li⁺ ions from transferring to dead corners, thus increasing the interface resistance and capacity attenuation. In addition, a large volume change in the electrode materials may also result in a loss of effective contact between the electrodes and electrolyte during repeated charging and discharging [48].

In TFLIBs, the appearance of interface resistance due to the intermediate layer formation is also a well-known phenomenon. One of the leading causes of interface resistance is the formation of space charge regions. Space charge regions are driven by high potential gradients, leading to Li-depletion near the cathode-SSEs interface in TFLIBs [49]. As a result of the difference in potential at such an interface, Li⁺ moves to a higher potential area, leading to Li-depletion and an increase in the interface resistance. In addition, the high interfacial resistance may also be a result of the interfacial chemical reactions that occur due to the mutual diffusion between the SSEs and electrodes [50]. These interfacial reactions irreversibly change the electrode material into inactive substances during operation. The inactive substances consume the Li⁺ but also hinder the Li⁺ migration between the electrolyte and electrodes, which significantly reduces the battery cycle life [51]. Reducing the solid interface impedance and improving the stability of the solid interface is the greatest challenge for improving the performance of TFLIBs. Recently, it has been reported that the solid interface can be modified by introducing an intermediate layer to reduce the interfacial resistance; this will be discussed in the next section.

2.2.3. Energy Density

The energy contained in batteries is a vital indicator of the battery’s performance, which directly determines the application fields. When evaluating the energy content of batteries, the specific energy (measured in Wh·kg⁻¹) and energy density (measured in Wh·L⁻¹) are commonly utilized [52]. TFLIBs exhibit a higher specific capacity, broader working voltage, and longer cycle life than traditional LIBs. However, due to the geometric structure and deposition technology limitations, meeting the energy and power density requirements is a challenging task for rapidly developing micro-medical devices, wireless communications, sensors, and other fields. Achieving high specific energy and energy density in electrode materials necessitates meeting two fundamental requirements: (1) High specific charge (Ah·kg⁻¹) and charge density (Ah·L⁻¹), indicating that the material contains a large amount of charge carriers per unit mass and volume; (2) A high (for the cathode) and low (for the anode) standard redox potential in their respective electrode redox reactions, which can achieve a high batteries voltage [17]. Additionally, both the positive and negative electrode reactions must be highly reversible to retain the specific charge during numerous cycles in rechargeable batteries. Furthermore, the energy density can be increased by designing thin film structures to improve the active material content and contact area of the electrode material; examples of this are 3D structuring or anode-free thin film Li-metal batteries.

2.2.4. Charging Rate

Recently, the fast charging of battery packs (for example, reaching 80% of the nominal capacity within 15 min) has become an important goal in developing electronic devices [53,54]. The ability to charge at a high rate allows for the rapid recharging of mobile electronic devices, which enhances the user-friendliness. The rate performance of TFLIBs depends mainly on the ions and electrons mobility of the electrodes, electrolyte, and across interfaces. In particular, the electrode’s poor ionic and electronic transport is considered a rate-determining step of modern batteries. Therefore, electrodes should meet two criteria for fast charging technology: high ionic and excellent electronic conductivity [55]. Numerous endeavors have been taken to enhance the rate performance of batteries by developing new electrode materials and modifying their structure. However, charging under a high load will accelerate degradation, which leads to a decline in the battery power and capacity. To accelerate the energy development of TFLIBs, it is critical to strike a balance between the rate capacity and energy density.

2.2.5. Manufacturing Cost

Another significant barrier to the widespread application of TFLIBs is the manufacturing cost, including the cost of the raw materials and the deposition processes. Despite the discovery of numerous high-performing electrode and electrolyte materials, high costs are incurred, particularly when preparing thin films. For example, when using PVD to deposit the NCM cathode, it would be costly to prepare the sputtering target. Additionally, one needs to maintain a low deposition rate to produce high-quality films during the deposition process, which is time-consuming. Therefore, it is crucial to lower the cost of producing TFLIBs. Lowering the manufacturing cost of TFLIBs can be achieved by optimizing the materials, designing novel architectures, and optimizing the deposition process.

3. Batteries Materials

TFLIBs can be assessed based on various factors, such as: (1) open-circuit voltage; (2) maximum working current; (3) discharge capacity; and (4) recyclability. Similarly to traditional LIBs, the open-circuit voltage is solely influenced by the electrode materials. The lifetime is determined by the solid interfaces that exist in the whole cell. The maximum working current that batteries can handle is highly dependent on the interfacial condition between the electrodes and SSEs. The electron and ion diffusivities in the materials are also important [17]. Therefore, exploring more promising materials can effectively improve the battery performance. This section reviews the recent developments in advanced materials for TELBs, including current collectors, electrodes, and electrolyte materials.

3.1. Current Collectors

The current collector (CC) not only serves as a conduit for the transportation of electrode materials, but also plays a crucial role in the collection and transportation of the electrons in TFLIBs [56,57]. It plays an essential role in battery’s (dis)charge performance. To guarantee excellent battery performance, CCs must satisfy the following requirements in TFLIBs: (1) High electrical conductivity to facilitate the efficient diffusion of the charge carriers within the electrodes [58]; (2) High robustness to improve the electrode stability during consistent charging and discharging procedures [59]; (3) High thermal conductivity to quickly disperse ohmic and chemical reaction heat within the batteries [60]; (4) High chemical inertness to ensure stability alongside the complex electrochemical reactions [61]; (5) Compact size and lightweight to infuse the batteries with high energy and power densities [62]; (6) Low cost to achieve the demands of mass production and commercialization [63].

In recent years, researchers have explored a variety of materials for use as CCs. CCs are typically metallic elements, such as Al, Cu, and Au [64]. In addition, inorganic compounds and metal alloys have been explored, such as TiN, Nickel alloys, stainless steel, etc. The range of the electrochemical stability voltages of CCs directly determines its application for either an anode or cathode. For example, Cu and Al are the standard CCs for anodes and cathodes. The electrochemical stability voltages of these CCs are shown in

Table 1. Various thin film current collectors for different electrodes.

Current Collectors	Electrode	Electrolyte	Voltage Window	Reference
Au	cathode	1 M LiClO4 in PC 1 M LiPF6 in EC/DMC	3–5 3–4.4	[65]
Ag	cathode	1 M LiClO4 in PC 1 M LiPF6 in EC/DMC	3–3.7	[66]
Al	cathode	1 M LiClO4 in PC 1 M LiPF6 in EC/DMC	1.5–5.5 1.5–5	[65]
Ni	cathode	1 M LiPF6 in EC/DMC	3–4.5	[67]
Stainless steel	cathode	1 M LiPF6 in EC/DMC	3–4.5	[68]
Stainless steel	cathode	1 M LiPF6 in EC/DMC	1.5–5.5	[69]
Cr	cathode/anode	1 M LiPF6 in EC/DMC	0–4	[70]
Ti	cathode/anode	1 M LiPF6 in EC/DMC	0–4	[70]
TiN	cathode/anode	1 M LiPF6 in EC/DMC	0–4.12	[69]
Carbon fiber paper	cathode/anode	1 M LiPF6 in EC/DMC	1.5–3	[71]
Stainless steel	cathode/anode	1 M LiPF6 in EC/DMC	2–3.4	[72]
Fe	anode	1 M LiPF6 in EC/DMC	0–3.2	[70]
Cu	anode	1 M LiPF6 in EC/DMC	0–3	[70]

. Although various CCs have been developed recently, they are challenging to apply on an industrial scale due to their poor mechanical strength and (electro)chemical stability. Poor adhesion between the surface of CCs and the electrode materials is also a significant contribution. The next generation of TFLIBs is aimed at achieving a higher capacity, longer lifespan, environmental protection, and economic price. This requires CCs to exhibit higher electrochemical stability, higher conductivity, lighter weight, and lower price. To enhance the performance of CCs, various approaches have been explored, including the design of an artificial layer on the CC’s surface, the optimization of materials, and the structure of CCs.

3.1.1. Surface Coatings

The coating is a simple method to enhance the conductivity and corrosion resistance of CCs to obtain a better performance. Surface coatings have been extensively studied to protect CCs from corrosion. Compared to adding a corrosion inhibitor to the electrolyte, coating layers can provide better protection because thicker layers can be formed on the current collector surface by controlling the deposition process and composition. However, the coating may result in higher contact resistance between the electrodes and the CC surface. The coating should inhibit corrosion and prevent the increase in contact resistance. Recently, many materials have served as a coating on CC surfaces to improve the performance, such as carbon [73], Ag [74], Ni [75], ZnO [76], and CuO [77]. Carbon materials are excellent coating layers; they enhance the rate performance, reduce the internal resistance, and maintain long-term cycling capacity retention, while also providing corrosion protection and significantly improving the thermal conductivity [73].

Large-area few-layer graphene films were also coated on an Al current collector by CVD to achieve a LIB with a higher rate performance and cycling stability [78]. LiCoO₂ (LCO) was papered as an active material to measure the electrochemical performance. It was found that three-layered graphene demonstrated the highest electrochemical performance and corrosion resistance, and the capacity retention is 82.9% after 100 cycles, which is a substantial increase over the Al current collector (75.8%). Graphene film was also coated on Cu current collectors via LPCVD. The graphene-Cu composite was demonstrated to effectively decrease the internal resistance, as well as enhance the Li₄Ti₅O₁₂ (LTO) anode material’s utilization efficiency. This improves the specific discharge capacity, rate capability, and cycle life, particularly at high current densities [79]. At a high C-rate (10 C), the LTO electrode achieved a capacity retention of 32%, with a modified current collector higher than LTO with bare Cu. Additionally, metals and metal oxides are good candidates to serve as coating layers.

3.1.2. Alloyed Current Collectors

In addition to the surface modification of the original collector, many new collector materials with excellent performance have been recently explored. The exploration of new-type current collector materials is mainly divided into three directions:

(1) Alloy CCs with transition metals, such as Ni, Cr, and Fe, to improve the mechanical strength and stability and to have a wider range of electrochemical stability voltages. TiN has emerged as a potential candidate for current collectors and can be produced through the magnetron sputtering of a Ti target under a nitrogen-argon atmosphere. Researchers have noted that TiN is a promising cathode current collector for high-voltage LIBs and has demonstrated superior oxidative stability in LiPF₆ and lithium bis(fluorosulfonyl)imide (LIFSI) electrolytes compared to other current collectors, such as Al or stainless steel [76,80].

(2) Enhance the ability of electrode materials to accommodate volume changes in order to maintain better contact with the electrode material, thereby improving the cyclability. TiNi alloys have recently received considerable attention because of their super-elastic and shape memory properties [81]. The super-elasticity of TiNi alloys is attributed to a stress-induced reversible martensitic transformation. A 50.3Ti–49.7Ni thin film prepared by DC sputtering was utilized as a current collector for Si film electrodes. The designed Si/TiNi electrode demonstrated a high initial columbic efficiency of 87%. Furthermore, the Si/TiNi film electrode achieve a high-capacity retention (46% for the first cycle) for up to 50 cycles, compared to Si film using a Cu current collector (28%). In addition, the electrochemical properties, cycle performance, and structural stability of the Si/TiNi electrode were significantly improved compared to Si electrodes using conventional Cu foil current collectors.

(3) Exploring low-density current collectors with high conductivity and a large surface area to improve the specific energy density. Carbon-based materials are of interest due to their lower density, superior flexibility, and chemical stability compared to metals. For this reason, several carbon-based CCs have been proposed, including reduced graphene-oxide films, carbon nanotube films, and graphene-coated films, to enhance the energy density and cycle stability [82]. Shichun Mu et al. developed a flexible, lightweight graphite film (GF) through the direct carbonization and graphitization of polyimide films (PI) (Figure 3a) [83]. The GF possesses an extraordinary planar conductivity ultrahigh electrical conductivity, of 1.07·10⁻⁶ S·m⁻¹, and low surface sheet resistance, resulting from its high planar oriented structure. It displayed an outstanding cycling performance (142.31 mAh·g⁻¹ at 2.0 C), with a capacity retention of 71.80% after 1000 cycles in the CSi half-cell, showing a 13.82% improvement over that using the Cu collector. In addition to maintaining the same stable cycling performance as metal current collectors, the extremely light weight of graphite films also resulted in a more than 25% improvement in the electrode gravimetric energy density (332 Wh·kg⁻¹) than that (260 Wh·kg⁻¹) of LiCoO₂||CSi full-cell with Cu and Al current collectors (Figure 3b).

3.1.3. 3D-Current Collectors

The surface of the CC exhibits a high specific surface area to enhance the adhesion of electrode materials to the CCs and increase the contact area. The CC’s surface should not be flat and should be modified to increase its surface area. To enlarge the surface area of CCs, surface treatment methods with different surface morphology have been developed. The 3D surface structures were fabricated using an anodization process, chemical etching, and microscale lithographic technology. The honeycomb-like surface of the aluminum current collector is etched by an anodization process in sulfuric acid and phosphoric acid mixed solution to achieve sufficient electrical contact between the electrodes and CCs [85]. It was shown that the contact between the cathode and Al CCs improved by 23% after anodization. After the surface treatment, the corrosion current peak reduces from 0.267 to 0.022 mA·cm⁻². The capacity retention of the LiCoO₂ electrode with an oxidation-treated aluminum current collector is 6.3% higher than with untreated aluminum foil at 5 C after 500 cycles. In addition, dendritic copper current collectors have been prepared by hydrogen evolution reactions (Figure 3c) [84]. The high surface structure of the dendritic copper collector promotes the migration of Li-ions, intending to inhibit the growth of Li dendrites. Compared to the planar copper foil, the dendrite copper foil exhibits low interface resistance and the stable Li-electrodeposition dendritic copper current collector showed low voltage hysteresis when driven stably without short circuit issues. A long cycle life of 500 h at 1 mAh·cm² was achieved in a symmetric performance, which resulted in a good cycling performance and cycle stability. A Li-metal anode with the cell also showed an excellent full-cell cycling performance during 40 cycles.

Chemical etching is also an effective treatment for current collectors to achieve a rough surface and better performance. Yang et al. designed vertically aligned CuO nanosheets on planar Cu CCs by simple wet chemical reaction (Figure 3d) [77]. The integration of the CuO array and Cu ensures good electron transfer. The vertical arrangement of channels between CuO nanosheets ensures the rapid diffusion of ions and reduces the local current density. The nanosheet structure of CuO remains unchanged and keeps good contact with the Cu collector during the cycle. An average CE of 94% can be maintained for 180 cycles in a VA-CuO-Cu||Li half-cell. The average CE of 92% can still be maintained with a CuO-Cu current collector with a high current density of 3 mA·cm⁻². An assembled full cell with a LiFePO₄ cathode and Li@VA-CuO-Cu as the anode delivered a capacity retention of 81.3% over 300 cycles.

In recent years, micro-scale lithography technology has been widely applied to produce three-dimensional surface structures. This technique can accurately control the surface morphology and etching rate. A surface-patterned cellular Cu@CuO composite current collector with a hexagonal blind hole array has been synthesized by lithography and solution immersion [86]. The results show that the batteries with the cellular Cu@CuO current collector maintain a reversible charge capacity of 354.1 mAh·g⁻¹ at 100 mA·g⁻¹, even after 200 cycles. Compared to the Cu-based CCs with a hexagonal blind hole array and bare Cu current collector, the new pattern gains a performance improvement in the reversible capacity, respectively, by 27.9% and 153.2%. The improved electrochemical property of the cellular Cu@CuO current collector is ascribed to the strengthened adhesion with the active materials, reduced contact resistance, and improved Li-ion diffusion capability. Fabricating a 3D collector is an efficient method to obtain excellent collector materials that can effectively improve the stability and enlarge the contact area between the current collector and the electrodes. This, in turn, improves the (dis)charge rate of the batteries. More importantly, the 3D collector can be applied to fabricate 3D batteries, which will significantly improve the energy density.

3.2. Anode Materials

The anode material influences the first cycle efficiency and cycling performance of TLIBs and directly determines the batteries’ capacity. The development goals of anode materials are an impressive specific capacity, excellent cycling stability, and low price. Based on the underlying reaction mechanism, anodes can be divided into three groups [87]: intercalation, alloying, and conversion materials.

(1) Intercalation anodes lead to an outstanding cycling ability as the volume change during the reaction is not apparent. However, their capacities are relatively low. The most typical intercalation anode material is carbon. Carbon materials combine good electrical conductivity with high crystallinity and stable charging and discharging behavior. It is currently the most commercialized anode material for LIBs.

(2) Alloying anodes can deliver a very large capacity, but the volume expansion during cycling limits their cycling ability. They mainly include metal anodes, such as Si, Ge, Sn etc.

(3) Conversion anodes mainly refer to transition metal oxides, sulfides, nitrides, phosphides, and fluorides, such as Fe, Mn, Co, Ni, Cr, and Cu. There is no place for Li-ion insertion and extraction in their spatial structure. Therefore, they do not conform to the traditional Li-ion insertion and extraction mechanism. The reaction with Li-metal at room temperature has been considered irreversible. However, conversion anodes provide specific capacities, around 1000 mAh·g⁻¹ or even larger, nearly twice that of graphite. In this part, we will review the most popular anode materials of recent years.

3.2.1. Li-Based Anodes

Li-metal has long been recognized as an optimal anode material for rechargeable batteries, owing to its exceptionally high theoretical capacity (3860 mAh·g⁻¹) and its extremely low negative electrochemical potential (−3.040 V vs. SHE) [88]. Unfortunately, Li-metal is prone to precipitate as dendrites during charging when used as an anode material. The formed Li-dendrites pose a significant risk of short circuits because they can penetrate the separator to cause short circuits. Moreover, isolated Li-metal will be produced during the discharging process, which inevitably reduces the performance [89]. The formed Li-dendrites cause many severe problems, mainly including (Figure 4a): (1) Cell short circuits; (2) Aggravation of adverse reactions; (3) Dead-Li growth; (4) Increased polarization; and (5) Large volume changes [90]. Many approaches have been put forward to make Li-metal anodes feasible in LIBs, including electrolyte additives [91], solid-state electrolytes [92], artificial solid electrolyte interphase (SEI) [93], and structured anodes [94]. In the past, many reviews have systematically summarized how to improve the Li-metal anode performance. In this section, the most recent progress in the surface modification of Li-metal and Li-alloyed anodes will be reviewed.

Surface Modification of Li-Anodes

Many methods have been developed for TFLIBs to prepare a robust artificial film before cell operation. These methods have been proven in practical approaches to improve SEI layers and suppress the Li-dendrites formation. Chemical and physical methods can be used to construct artificial SEI layers. The chemical approaches mainly include gas and liquid treatments. Koch et al. have shown that a passivation layer Li₃N can be deposited on the Li-metal surface by reacting with N₂ at room temperature. The formed Li₃N layer exhibits better stability and forms good protection for Li-metal [96]. In addition to gas treatment, liquid treatment is another effective method to coat a protective film on the surface of the Li-anode. Li et al. prepared an artificial SEI layer with a uniform distribution of Li₃PO₄ through an in-situ reaction of polyphosphoric acid with Li-metal and its native film [97]. The Li₃PO₄ layer can promote the Li-ions’ transportation between the Li-metal anode and SSE’s interface and induce a smooth surface with a high Young’s modulus. This is sufficient to hinder the growth of Li-dendrites by mechanical protection.

Advanced thin film manufacturing methods, such as PVD, CVD, and ALD, can minimize the corrosion of Li-metal in the air during processing and meet the challenge of the low melting point of Li-metal. This is promising to control the coating thickness, uniformity, consistency, and defect density more accurately. Many inorganic materials with strong mechanical strength have been deposited as artificial SEI films, such as carbon [98], Al₂O₃ [93], BN [99], and LiF [100]. Carbon is the most common inorganic material used as artificial SEI films, which result in the formation of a stable and durable SEI film, which effectively reduces the growth of the Li-dendrites [98]. Al₂O₃ films have often been fabricated by spin-coating or ALD. The Al₂O₃ films prepared through different physical approaches possess different structures and morphologies and have demonstrated the effective suppression of Li-dendrite growth [101]. Applying chemical vapor deposition provides another effective way to improve the interface stability of Li-metal anodes. 2D boron nitride (BN) coated on the surface of the Li-metal anode is an excellent interfacial protection layer. Dendrite-free Li-morphology is realized because its flexibility, chemical stability, and mechanical strength come from its strong interlayer bonds and ultra-thin thickness [99]. Thus, physical coating represents a straightforward and effective approach to protecting Li-metal anodes.

Li-Alloyed Anodes

It has been discovered that using inorganic materials that react with Li-metal anodes to form alloys significantly reduces the interfacial resistance. Furthermore, ultra-thin and conformal alloy coatings can enhance the solid electrolyte wettability to Li-metal. It eliminates the separation between the solid electrolytes and Li-metal [102]. The alloying reactions of Li-metal with metallic elements, such as Mg [103], Al [104], Ge [105], Sn [106], Ag [107], and Au [108], or semi-metallic elements such as C [98] and Si [104], have been explored in recent decades. Al-coated garnet deposited by electron-beam evaporation exhibits excellent wettability to Li-metal through the reaction to form an Li_xAl alloy, resulting in reduced interfacial resistance [104]. A thin Au layer is sputtered onto the polished surface of SSEs to flatten the interface further, reducing the resistance and preventing short-circuiting [108]. In addition to metallic elements, some nonmetallic elements can also react with Li-alloys. For example, a 10 nm Si coating is deposited on a garnet electrolyte by PECVD. The silicon layer is in perfect contact with the garnet and reacts with the Li-metal to create a lithiumphilic surface, which significantly reduces the interfacial resistance between the garnet electrolyte and Li-metal anode [104]. The alloying interfacial layers can effectively decrease the interfacial resistance and block the side reaction between the SSEs and Li-metal. In addition, many compounds also exhibit a high reactivity with Li-metal anode. For example, with the ultrathin, conformal ZnO film coatings, ZnO can react with Li-metal during cycling to create better contact with the garnet electrolyte surface, which leads to a significant reduction in the interface resistance [13].

3.2.2. Si-Based Anodes

Pure Si Anodes

Pure amorphous and crystalline Si both serve as negative electrode materials. However, it is reported that the crystal structure is inactive before it becomes amorphous, followed by the initial alloying with Li-ion. Obrovac et al. [109] carried out one of the earliest clarification works. They used XRD to investigate c-Si and observe its conversion to partially lithiated a-Si in the initial lithiation process. The c-Si was transformed into partially lithiated a-Si. Then a-Si crystallized to Li₁₅Si₄, and the fully lithiated material existed at room temperature when the voltage decreased to 50 mV. The formed phase is known to cause high internal stresses, particle cracking, poor electrical contact, and capacity fading. However, numerous experiments and calculation results have demonstrated that the cycling performance of the Si film electrode can be improved by reducing its thickness [110,111,112]. Ohara et al. found the best performance of a Si thin film when the thickness was 50 nm, which maintained a high specific capacity of approximately 3100 mAh·g⁻¹ and a high coulombic efficiency of 100% at 12 C after 1000 cycles [113]. This was early research on the electrochemical performance of Si thin film anodes. However, there are still some problems with pure silicon thin film anodes. The main research question is how to manage and limit the effects of the total volume change. In addition, a highly lithiated phase forms between the Si electrode and Cu CCs at the onset of the lithiation process. The fast-diffusion paths and non-lithiated Si regions are both being filled with Li⁺ ions at the same time. (Figure 4b). However, pure Si films exhibit poor intrinsic electronic conductivity of 10⁻³ S·cm⁻¹ and a low Li-ion diffusion rate (diffusion coefficient between 10⁻¹⁴ and 10⁻¹³ cm²·s⁻¹). This limits the Li-ion diffusion rate in the whole process. Therefore, it is also important to improve the effective conductivity of silicon layers to improve the power performance of TFLIBs.

In recent years, many proposed solutions and strategies have been put forward for controlling and buffering the problems of the high-volume expansion of silicon anodes, as shown in Figure 5a. Modifying the substrate surface is very important to obtain long-term cycle stability for a pure Si thin film system. For example, the copper foil is roughened with sandpaper to enlarge the contact area with the deposited amorphous Si film and help the film with a higher tolerance with the volume changes [114]. In addition, many chemical etching methods have been used to modify the surface of silicon anodes. A substrate was etched in FeCl₃, which can accommodate volume changes and enhance the capacity of the Si anodes [115]. Another original method is to treat the substrate through plasma-immersion-lanthanum-ion implantation so that the capacity retention rate of the silicon film anode can reach 93% in the 80th cycle [110]. Electrodeposition, also used to obtain copper foil with a rough surface, has been utilized to enhance the performance of a-Si thin film anodes. Unfortunately, it was demonstrated that the capacity retention of the anode is only 42% after the first discharge [116].

Although the above methods can effectively modify the surface of silicon films, their shape and etching thickness cannot be accurately controlled. Therefore, patterned films obtained by lithography are an excellent surface modification strategy to accommodate the volumetric changes. The electrochemical characteristics of Si film anodes have been improved by designing a variety of novel patterns and structures. For example, lozenge-shaped tiles (Figure 5b) [118] are produced by covering the mask on the substrate surface during deposition. As a result of the good stress suppression effect between the Si tiles, the good adhesion, and the improved mechanical stability of the patterned Si anodes during (de)lithiation, a high capacity retention of 91% can be achieved after 100 cycles. The study and comparison of planar and shaped Si nanostructures clearly show the advantages of the 3D patterned electrodes, including the high specific surface area, increased current density, and improved mechanical stability. In our group, we have designed an n-type doped a-Si honeycomb structure utilizing lithography (Figure 5c). The voltage distribution of the Si anode is displayed in Figure 5d, with inserted SEM images that correspond to different Li-contents. The mechanical deformation of the honeycomb wall is caused by the internal stress generated in the lithiation process [119]; 3D patterned Si anodes can accommodate the volumetric changes effectively to maintain high capacity and improve the energy density of TFLIBs. However, there are still many aspects to be improved, including the design of the pattern shapes and the optimization of the size, which will be an excellent opportunity to design even better pure-silicon thin film anodes.

Composition modification is also an excellent promising approach to enhance the performance of Si electrodes. This could significantly reduce the expansion and contraction of Si anodes during (dis)charging and improve the conductivity. In recent years, there have been numerous investigations into combining active Si anodes with other elements, such as Cu [120], Mo [121], Fe [122], Al [123], B [124], Y [125], and Ti [126]. The formed metal silicide (Si-Mo, Si-Cu, Si-B, Si-Ti) films reduce the resistance, and improve the mechanical properties due to the additional chemical bonds.

It was also found that a SiN layer coated on a Si surface can effectively enhance the electrochemical performance of Si anodes. Forming a Li₃N film at the surface improves both the storage capacity and ionic conductivity. In addition, a thinner SiNx layer was more effective because SiN prevents the transport of Li⁺ and the decline in capacity [127].

A conductive and mechanical buffer multilayered structure of Si anodes with interlayers of Fe, Al, Y, and Ti can accommodate volumetric changes. It is favorable for the loner cycle life of the electrodes. Sputtered Au-Ti/Si/LIPON/SnO₂ multilayered anodes with a LiPON/SnO₂ layer acting as a porous SEI layer showed no visible capacity loss over 200 cycles [128]. However, the challenge with multilayered Si-based anodes is optimizing the thickness, density, and quantity of layers. A sufficient amount of active Si must be available, while keeping a good Li⁺ diffusion rate and high CE. Although the capacity of silicon alloys and silicon composite is slightly lower than that of pure silicon, it is still higher than most other anode materials. Applying such materials can effectively enhance the ion conductivity of Si and reduce the influence of volume changes, thus improving the cycling performance.

Recently, Si oxide (SiO_x, 0 ≤ x ≤ 2) has attracted much attention in developing high-performance LIBs because of its ultrahigh theoretical Li-storage capacity, low working potential, earth-abundancy, and good cycle life [129]. Si-oxide materials, including SiO, SiO₂, and nonstoichiometric SiO_x, have various electrochemical properties. The theoretical capacities of SiO and SiO₂ are 2600 and 1965 mAh·g⁻¹, respectively [130,131]. The Li-storage capacity decreases gradually with the oxygen content in SiO_x. In addition, the volume changes of SiO and SiO₂ are smaller than that of pure silicon during (dis)charging. The SiO anode shows a small volume change of 150% after lithiation. The volume change is decreased from 300% for Si to 100% for SiO₂. With the increase in the oxygen content in silicon-based materials, the volume changes during (dis)charging decreases, and the cycle stability improves gradually.

3.2.3. Metal Oxides Anodes

Metal oxides are of great interest for TFLIBs because they can store Li-ions [132]. In addition, some metal oxides, such as iron and manganese, are naturally abundant and, thus, low-cost. The size, morphology, and orientation of metal oxide nanostructures can also be easily adjusted. This makes it possible to systematically investigate the relationship between their structure and electrochemical properties. Metal oxides can be reversibly reduced and oxidized by the following general equation:

$${\mathrm{M}}_x{\mathrm{O}}_y+2\mathrm{y}{\mathrm{L}\mathrm{i}}^{+}+2\mathrm{y}{\mathrm{e}}^{-}\to \mathrm{x}\mathrm{M}+\mathrm{y}{\mathrm{L}\mathrm{i}}_2\mathrm{O}$$

(1)

According to the crystal structure of metal oxides, the electrochemical mechanisms of these anode materials can be classified into three types [133]: (i) Intercalation reaction: intercalated compounds, such as Li₃VO₄, LiV₃O₈, Li₄Ti₅O₁₂, and TiO₂; (ii) Alloy reaction: oxides, such as SiO₂, SnO₂, GeO₂, and their mixed oxide derivatives; (iii) Conversion reaction: metal oxides, such as Co₃O₄, MoO₂, Fe₂O₃, V₂O_5, and NiO. In recent years, various oxide-thin film anodes with high electrochemical performance have been synthesized for TFLIBs (

Table 2. Various anode materials for TFLIBs.

Metal Oxides Anode	Reaction Mechanism	Initial Discharge Capacity (mAh·g−1)	Coulomb Efficiency	Reference
LTO	intercalation	313	95.0%	[136]
TiNb2O7(TNO)	intercalation	398	92.0%	[137]
TiNb2O7	intercalation	460	99.0%	[138]
SrLi2Ti6O14	intercalation	175	91.4%	[139]
SnO2	alloy	650	99.5%	[140]
LNVO	conversion	871	98.5%	[141]
NiFe2O4	conversion	849	67.7%	[142]
Co2O3	conversion	1235	56.0%	[143]
Fe2O3	conversion	951	100%	[144]
MoO2	conversion	1510	-	[145]
ZnO	conversion	1200	-	[146]
MoOx	conversion	1118	81.8%	[147]
RuO2	conversion	1200	33.0%	[148]
ZnOS	conversion	1271	64.1%	[149]
In2O3	conversion	1083	81.5%	[150]

). However, using metal oxides also brings several unpleasant difficulties, such as poor electrical conductivity, poor ion transportation [134], and severe volume changes [135] over the (dis)charging process. All of these factors result in the pulverization of thin film electrodes. Further efforts should be made to develop thin film metal oxide anodes to address these concerns and exploit high-efficiency electrochemical energy-storage devices.

3.3. Cathode Materials

The cathode films used in the early TFLIBs were mainly composed of binary compounds without Li-element, such as TiS₂ [151], MoO₃ [152], and V₂O₅ [153]. The preparation conditions are relatively simple, but the cycle performance is poor. They are classified as first-generation cathode films. Later, intercalation compounds were developed, such as LiCoO₂ [154], spinel LiMn₂O₄ [155], and olivine LiFePO₄ [156]. They have a higher voltage and show good cycling behavior. The application of these materials in TFLIBs is rapidly growing at present. The cathode films made of intercalation compounds are called second-generation cathode films. LiCoO₂ is widely applied in mobile electronic devices as a commercial cathode material. However, further advancement is necessary to meet the pressing demand for energy storage systems featuring high energy density and rapid charging rates. The industry requires electrode materials that demonstrate high voltage, specific capacity, and exceptional rate performance. This should be achieved through doping, exploration, and structural adjustment in the second-generation cathode materials.

3.3.1. High Voltage Cathodes

In recent years, the energy density of LIBs has been significantly improved, mainly through boosting the batteries’ materials and optimally engineering the design. The energy density of LIBs is calculated by integrating the operating voltage with respect to the storage capacity. As a result, enhancing the working voltage has become a critical approach to boosting the energy density of TFLIBs. As the anode materials work at low potentials near to Li-metal, increasing the total battery voltage highly depends on the cathode materials [157]. Dependent on the voltage range, three classes of materials can be considered for cathode host materials: (1) the 3-V cathodes, e.g., MnO₂, LiFePO₄; (2) the 4-V cathodes, e.g., LiCoO₂, LiCo_1−y−zNi_yAl_zO₂, LiNi_1/3Mn_1/3Co_1/3O₂, LiMn₂O₄; (3) and the 5-V cathodes such as spinel LiNi_0.5Mn_1.5O₄ (LMNO) and its related doped structures, olivine LiCoPO₄, inverse spinel LiNiVO₄, and fluorophosphate Li₂CoPO₄F [158]. Only when 5 V cathode materials are used, the energy density of Li-ion batteries can reach the industrial requirement of 500 Wh·kg⁻¹. This section reviews the recent development of LMNO and LiCoPO₄, the most commonly used high-voltage cathode materials.

Spinel LiNi_0.5Mn_1.5O₄ (LMNO) is considered to be an excellent cathode for next-generation TFLIBs for EV and grid-scale electric device applications because of its impressive working voltage of ~4.7 V (vs. Li⁺/Li), low price, long cycle life, and excellent rate performance [159]. However, several issues and challenges continue to be significant in the development of LMNO: (1) the unstable SEI formation at high operating voltages; (2) the dissolution of Mn during electrochemical cycling; and (3) the structural transition from the LiM₂O₄-type spinel structure to M₃O₄-type spinel, and rock-salt structures at the surfaces [160]. Numerous endeavors have been undertaken to enhance the electrochemical characteristics of LMNO, such as doping, controlling morphology, particle size, the degree of cation ordering, and surface modification [161].

Among these solutions, surface modification is the most efficient and simple way to prevent harmful side reactions upon the electrochemical cycling of thin-film LMNO, which enhances the overall capacity and cycle life. An efficient surface passivation layer prevents the direct contact of the active particles with the electrolyte. It suppresses the phase transition and side reactions at the particle surface, thus significantly enhancing the electrochemical performance of the cathode materials. The coating compound materials mainly include metal oxides (Al₂O₃ [162], TiO₂ [163], ZnO [164]), metal fluorides (MgF₂ [165], AlF₃ [166], CeF₂ [167]), polyanionic compounds (Li₃PO₄ [168], LiFePO₄ [169], Li₂SiO₃ [170], LiPON [171]), and carbon-based materials (carbon [172], graphene oxide [173]). The coating layers minimize the harmful side reactions between LNMO and the electrolyte to enhance the reversible (de)intercalation of Li-ions in the (dis)charge process. For example, an Al₂O₃-modified LMNO thin-film electrode prepared by PLD exhibits excellent (dis)charge performances when elevated at room temperatures. It delivers a discharge capacity of 126.3 mAh·g⁻¹ after 100 cycles at 25 °C, corresponding to 95.4% of its initial value. Even at 55 °C, 84.1% of its initial discharge capacity can be maintained at the 100th cycle, and the rate capability has also been improved substantially [162]. The coating can also increase the ionic conductivity of LMNO. Lv et al. [171] coated an ~15 nm amorphous LiPON thin film onto an LNMO thin film electrode by magnetron sputtering. It was shown that the LNMO/LiPON electrode exhibited stable reversible capacities (higher than 105 mAh/g) following cycling to 4.9 V after 150 cycles. LiPON serves as the artificial interface film, allowing Li-ions to diffuse through and suppress the formation of a thick insulating surface. EIS studies of such electrodes show lower interface and charge transfer resistances with lower electrode polarization. The uniformly distributed LiPON film on the LNMO cathode surface can separate the direct contact between the LNMO and electrolyte, inhibit the side reactions of Mn dissolution, and provide a stable interface for ionic transport. The LiPON coating on LNMO formed a space charge layer, maintaining a slightly higher manganese ion valence state.

Doping is also an excellent approach to enhance the Li-ion conductivity and structural stability by introducing different ions in the cathode host materials [174], mainly including metal cations such as Al³⁺ [175], Cr³⁺ [176], Co³⁺ [177], Fe³⁺ [177], and Ru⁴⁺ [178]. It was found that doping different ions will have different effects. For example, Mg-doped LNMO is reported to possess lower polarization, enhance Li-ion conductivity at room temperature, and induce good stability upon cycling, even at high C-rates. Cr-doped material results in better chemical and structural stability with less Jahn–Teller distortion for the Cr-doped spinel materials. Due to it having a smaller lattice parameter, the co-doped LNMO has more improved electrochemical properties than LNMO. Fe doping can promote the enrichment of specific surfaces and suppress the formation of thick SEI films, significantly enhancing the electrochemical performance of 5 V spinel cathode materials. In addition to doping metal ions, the performance of LNMO can also be improved by substituting anions for O. It was shown that fluorine doping can improve the electrochemical properties. It can suppress the formation of NiO impurity and simultaneously reduce the voltage polarization [179].

Olivine-structured LiCoPO₄ (LCP) has attracted considerable interest as a potential cathode material for LIBs due to its high redox potential (4.8 V vs. Li⁺/Li) and theoretical capacity (167 mAh·g⁻¹) [180]. However, LCP exhibits some unsatisfactory characteristics, including low specific capacity, poor cycle stability, and poor rate performance. Through surface modification, ion doping, particle size reduction, and morphology control [181,182,183], great efforts have been taken to enhance the electronic and ionic conductivity of LiCoPO₄. Surface modification is a common strategy to enhance the cycling stability of the LCP cathode by protecting the LCP from direct contact with electrolytes. This inhibits surface degradation caused by the HF attack induced by electrolyte decomposition [184]. Many materials have been used to modify the surface of LCP to enhance its electronic conductivity and electrode stability; examples of this are carbon [185], NiO [186], and AlF₃ [187]. An LCP was coated with an 8–10 nm continuous NiO layer, which can reduce the volumetric stresses and strains that occur during Li⁺ insertion and extraction.

3.3.2. High-Capacity Cathodes

Cathode materials essentially limit the battery’s storage capacity because many high-capacity anode materials already exist, such as Li-metal and Si. To improve the output capacity of TFLIBs, it is therefore essential to explore new promising high-capacity cathode materials. Recently, many novel crystal structure cathodes have been reported, as shown in Figure 6 [188]. Layered transition metal (TM) oxides are considered very promising cathode materials, with high energy density, owing to their high theoretical capacity (~280 mAh·g⁻¹) compared with that of olivine (~170 mAh·g⁻¹) or spinel (~150 mAh·g⁻¹) [189,190]. The most common layered cathode material is LiCoO₂ (140 mAh·g⁻¹). However, an LiCoO₂-layered oxide can only provide a limited practical capacity due to its narrow potential window (~4.3 V vs. Li/Li⁺). Recently, Li- and Ni-rich layered oxide materials have received more attention compared to other layered cathodes, owing to their high capacities. The capacities of Li-rich and Ni-rich layered materials are between approximately 1.25 and 1.5 times higher than that of the relatively low Ni content layered materials, such as the commercialized LiNi_1/3Co_1/3Mn_1/3O₂ (160 mAh·g⁻¹) [188]. In recent years, many high-capacity thin-film cathode materials have been explored. An overview is given in

Table 3. Various cathode materials for TFLIBs.

Metal Oxides Anode	Crystal Structure	Capacity	Voltage Range (V)	Rate Performance	Reference
V2O5	Layer	142 mAh·g−1	2.5–4.0	86.7 mAh·g−1 at 56 C	[191]
LiV2O5		36 μAh·cm−2·μm−1	2.4–3.6		[192]
HT-LiCoO2	Layer	135 mAh·g−1	3.0–4.2	103 mAh·g−1 at 10 C	[193]
LiNiO2	Layer	23 μAh·cm−2·μm−1	2.8–4.2		[194]
LiNi0.8Co0.2O2	Layer	125 mAh·g−1	2.8–4.4		[195]
Li2MnO3	Layer	267 mAh·g−1	2.0–4.8	147.84 mAh·g−1 at 9.3 C	[196]
LiNi0.8Co0.15Al0.05O2 (NCA)	Layer	197 mAh·g−1	2.5–4.2		[197]
LiNi1/3Co1/3Mn1/3O2 (NMC333)	Layer	177 mAh·g−1	2.8–4.5	60 mAh·g−1 at 2 C	[198]
LiNi0.5Mn0.3Co0.2O2 (NMC532)	Layer	167 mAh·g−1	3.0–4.5		[199]
LiNi0.6Mn0.2Co0.2O2 (NMC622)	Layer	79 mAh·g−1	2.5–4.5		[200]
Li1.2Mn0.55Ni0.15Co0.1O2	Layer	250 mAh·g−1	2.0–4.8		[201]
LiMn2O4 (LMO)	Spinel	82 mAh·g−1		67 mAh·g−1 at 348 C	[202]
Li 0.92Co0.65Mn1.35O4	Spinel	340 mAh·g−1	1.6–4.5	268 mAh·g−1 at 10 C	[203]
LiCoMnO4	Spinel	107 mAh·g−1	3.0–5.0		[204]
LiNi0.5Mn1.5O4	Spinel	117 mAh·g−1	3.0–5.0	108 mAh·g−1 at 10 C	[205]
LiMnPO4	Olivine	24 mAh·g−1	3.5–4.4		[206]
LiFePO4	Olivine	159 mAh·g−1	3.0–4.0	153 mAh·g−1 at 7.5 C	[207]

.

Li-rich oxides (Figure 6a) possess higher specific capacities (>250 mAh·g⁻¹) and energy densities (>900 Wh·kg⁻¹), which meet the requirements of the next-generation rechargeable TFLIBs [208]. The high capacity is provided by the accumulated cation and anion redox reactions of the Li-rich oxides. The high manganese content ensures good safety and low cost. Li-rich oxides are a group of lithium transition metal oxides (Li_1+xTM _1−xO₂) with a Li/TM ratio greater than 1 (0 < x < 1), such as Li₂MnO₃ (200 mAh·g⁻¹) [209], Li₂RuO₃ (270 mAh·g⁻¹) [210] and Li₃IrO₄ (340 mAh·g⁻¹) [211]. In recent years, a lot of novel materials with high capacity have been designed and synthesized by partially replacing the Li and TM elements in Li-rich oxides, such as Li_1.9Cr_0.3Mn_0.8O₃ (250 mAh·g⁻¹) [212] and Li_1.144Ni_0.136Mn_0.544Co_0.136O₂ (301 mAh·g⁻¹) [213]. However, layered Li-rich cathode materials still face severe problems, such as a low initial coulombic efficiency, severe voltage decay, significant capacity losses, poor rate capability, and transitional metals dissolution [214]. In recent years, a variety of approaches have been explored for improving the electrochemical performance of Li-rich oxide cathodes. Substitution/doping and coating are two of the main strategies to improve the electrochemical performance. Different elements and radicals, such as Mg [215], Al [216], Ti [217], Zn [218], Sn [219], Y [220], B [221], F [222], and Cl [223], have been introduced into the framework of layered Li-rich oxides. Both TM-site substitution and surface substitution/doping improve the initial CE, rate capability, and cycling stability. For example, substitution/doping with ions that possess a large ionic radius (e.g., K⁺) or high electronegativity (e.g., F⁻) can hamper the migration of the TM ions, and thus mitigate the voltage decline. In addition to doping, a great variety of materials have been selected to modify the surface of layered Li-rich oxides, such as oxides [105], fluorides [224], phosphate [225], and silicates [226]. The main function of such protective coatings is to reduce the direct contact between the cathode and electrolytes, thus reducing the parasitic electrolyte oxidation and TM dissolution. It has also been observed that the coatings in the layered Li-rich oxides can delay the structural changes near the surface. As a result, such coatings usually result in longer cycle life materials.

Ni-rich layered oxide is another interesting cathode material for high-capacity TFLIBs, and has attracted significant attention. LiNiO₂ is the most typical representative of Ni-rich layer oxides. It shows a higher discharge capacity of >220 mA·h g⁻¹ and a lower cost than LiCoO₂ [227]. However, the difficulty in synthesis, serious Li/Ni cationic mixing, and poor thermal stability of LiNiO₂ hinder its commercialization. Many transition-metal dopings have been developed to partially replace Ni and resolve the problems of LiNiO₂ [228]. The most successful solutions are ternary layered oxides, LiNi_xCo_yMn_1−x−yO₂ (NCM) [229], or LiNi_xCo_yAl_1−x−yO₂ (NCA) [230]. They generally exist in particles and are essential as high-performance cathode materials. During the charging (oxidation) process, the oxidation of NCM starts with Ni²⁺ to Ni³⁺, followed by Ni³⁺ to Ni⁴⁺, and finally, Co³⁺ to Co⁴⁺. For Mn⁴⁺, it remains electrochemically inactive. In general, the oxidation of Ni²⁺ to Ni³⁺/Ni⁴⁺ occurs at relatively low operating potentials (<4.3 V vs. Li⁺/Li). The oxidation of Co³⁺ to Co⁴⁺ only appears above 4.3 V [231]. As the applied upper cut-off voltage is often below 4.3 V, Ni-rich layered oxides’ capacity mainly originates from Ni’s redox chemistry [232]. Therefore, by increasing the Ni content and reducing the Co and Mn content, Ni-rich layered oxides provide higher capacity (e.g., LiNi_0.6Mn_0.2Co_0.2O₂ (NCM622): ∼180 mAh·g^–1, LiNi_0.8Mn_0.1Co_0.1O₂ (NCM811): ∼200 mAh·g^–1). However, this occurs at the expense of the capacity retention and thermal stability [233].

Although these Ni-rich layered oxides exhibit high capacities, many problems and challenges hinder their practical applications. It can be summarized as follows [234]: (1) Residual Li-compounds. The existence of residual Li compounds leads to the gelation of polyvinylidene difluoride binders during electrode preparation; (2) Microcracking of particles. During the (dis)charging process, the microcracks are caused by anisotropic lattice distortion and mechanical strain related to the different phase transitions, especially in Ni-rich layered oxides with more than 80% nickel content; (3) Rock-salt phase formation. The electrolyte infiltrates the microcracks and forms an excessive interface layer between the positive electrode and electrolyte. These problems result in the loss of the active transition metals (decreased capacity and working voltage) and the migration of the transition-metal ions to the anode. This affects the optimized capacity ratio of the cathode and anodes and increases the impedance of the solid electrolyte interphase.

To date, many methods have been used to solve the issues of Ni-rich layered oxides and improve their electrochemical performance, including primary particle engineering, surface coatings, doping, and concentration gradient designs. For thin-film battery systems, surface coatings are a simple and effective method. Introducing coating materials onto the surface of Ni-rich layered oxides avoids direct contact with the electrolyte, thus minimizing the parasitic reactions. It also sets a kinetic barrier to O₂ evolution. Different types of surface coating materials, such as oxides (Al₂O₃, ZrO₂, TiO₂, and Co₃O₄), fluorides (AlF₃), and phosphates (Li₃PO₄), have been reported [235].

Moreover, the electrochemical inactive coating materials inevitably reduce Ni-rich layered oxides’ capacity and energy density. In addition, advanced technologies to deposit these coating materials also need to be improved to enhance the quality of the coatings. Similarly to surface coatings, bulk doping is also an effective and straightforward approach for regulating the structure of Ni-rich layered oxides and, thus, their electrochemical performance. To date, doping with a single cation, doping with a single anion, and co-doping with multiple ions have been reported to modify Ni-rich layered oxides. However, this is mainly conducted for powder batteries. It is more complicated to apply these approaches in thin-film batteries.

3.4. Electrolytes

The electrolyte is sandwiched between the positive and negative electrodes, which is critical for the optimal electrochemical stability of batteries during long-term operation. A good performance requires an impressive ionic conductivity, Li⁺ transference number, and electrochemical and thermal stability. Significant mechanical strength is also desired. Until recently, the absence of liquid electrolytes working with electrochemical windows above 5 V limited the application of electrode materials. However, the recent breakthroughs in exploring inorganic solid electrolytes have opened up new possibilities for using alternative electrode materials. Excellent inorganic solid electrolytes can help gain thermal and chemical stability, impressive ionic conductivity, and prolonged cycle life [236]. The preparation method, however, strongly influences the performance of solid electrolyte films.

Moreover, the deposition rate of various elements may vary considerably. However, a stable ion-conducting structure requires various elements, generating some deposition challenges. Solid electrolytes can be divided into three development stages, according to the complexity of preparation [17,237]. The early solid electrolytes were clearly limited by the preparation methods, mostly binary compounds and their composites [238], such as Li₃N, Li₂S, and Li₂S-SiS₂-P₂S₅. Although solid-state electrolyte systems, such as GeS₂, exhibit a conductivity higher than 10⁻⁵ S·cm⁻¹, they are highly reactive in the humid air. Subsequently, more stable oxygen-containing inorganic salts have been developed; examples are amorphous borate (Li₂O-B₂O₃-SiO₂), silicate (Li₂O-V₂O₅-SiO₂), and other systems called first-generation solid-state electrolytes. Although these systems are relatively easy to synthesize, they are all highly alkaline, absorb water easily and are challenging to operate. In addition, they have a low ion conductivity, which limits the development of TFLIBs.

With the development of new reactive deposition technologies, a second-generation solid electrolyte film, based on LiPON, with excellent performance has been developed [17]. TFLIBs have made considerable progress, and various battery systems and structures have been developed. With these new thin film preparation methods, which facilitate greater control of the thin-film components, such as PLD and ALD, many thin-film oxide-based electrolytes with excellent performance have gradually been obtained. These inorganic film systems can mainly be divided into four structural types: (1) Perovskite structure: LiO_0.33La_0.56TiO₃ (Li₃xLa_2/3−xTiO₃, x = 0.1–1) with high ionic conductivity; (2) NASICON structure: Li_1+xA_xTi_2−x(PO₄)₃ (LATP) with high chemical compatibility with Li-metal [239]; (3) LISICON and Thio-LISICON structure; (4) Garnet structure: Li₇La₃Zr₂O₁₂ (LLZO) with higher environmental stability [240]; (5) Sulfides: Li₁₀GeP₂S₁₂ and Li_9.54Si _1.74P_1.44S_11.7Cl_0.3 with high conductivity 1.2·10⁻² and 2.5·10⁻² S·cm⁻¹ [241]. The crystal structures of these inorganic solid electrolytes are shown in Figure 7, which are regarded as third-generation solid electrolyte membrane systems.

3.4.1. Perovskites

In 1971, Takahashi et al. [242] first reported perovskite-type solid electrolytes with the common formula ABO₃ (A=La, Sr, or Ca; B=Al or Ti). Perovskite has a face-centered cubic structure (Figure 7a), where A is at the top corner of the cube, B is at the body center, and O is located at the face center [243]. The perovskite-type Li_2/3−xLa_3xTiO₃ (LLT) has been extensively studied in TFLIBs. It is a promising electrolyte and one of the most conductive oxide electrolytes available because of its impressive crystalline conductivity at room temperature. However, LLT is plagued by high grain boundary resistance, significant electronic conductivity, and a tendency to react with Li-metal anode (reduction in the Ti⁴⁺) [244]. LLT has been successfully deposited on MgO and SrTiO₃ substrates by PLD and reactive sputtering [245,246]. In 2002, Ohta et al. deposited a predominantly single-crystalline LLT film onto (001)-oriented SrTiO₃. At 190 °C, the deposited LLT film demonstrated a high conductivity of 2.5·10⁻² S·cm⁻¹, with an activation energy of 0.37 eV, which is similar to the bulk LLT. The extrapolated results suggest that it has a conductivity of 1.5·10⁻⁴ S·cm⁻¹ at 25 °C. Moreover, PLD has also been employed to deposit amorphous LLT films [247,248]. Ahn and Yoon optimized the thermal processing conditions to achieve maximum conductivity while maintaining its amorphous structure [247]. The ion conductivities of the deposited films at 500 °C exceeded 1·10⁻⁵ S·cm⁻¹ measured at 25 °C. However, when LLT is deposited directly onto LCO at a deposition temperature of over 400 °C, less Li-ion capacities are obtained galvanostatically. This is due to the interfacial reactions and inter diffusion that occur when deposited at a high temperature. However, rapid thermal annealing can induce the crystallization of the LCO film and maintain its amorphous structure, which can effectively increase its capacity [248]. Other perovskite-structured solid electrolytes have also been used in TFLIBs, see

Table 4. Various solid electrolyte materials for TFLIBs.

Electrolyte	Crystal Structure	Deposition Method	Ea (eV)	Ionic Conductivity (S·cm−1)	Reference
LiPON		PLD		1.5·10−8	[249]
LixNbyO		ALD	0.62	6.4·10−8	[250]
Li0.33La0.56TiO3 (LLTO)	Perovskite	PLD	0.35	3.5·10−5	[251]
epitaxial Li0.33La0.56TiO3	Perovskite	PLD	0.34	6.7·10−4	[252]
Li0.17La0.61TiO3	Perovskite	PLD	0.25	3.8·10−4	[253]
Li0.5La0.5TiO2	Perovskite	PLD		1.1·10−5	[248]
Li3OCl	anti-perovskite	PLD	0.36	8.9·10−6	[254]
Li 1+xAlxGe2−x(PO4)3 (LAGP)	NASICON	sputtering		>10−4	[255]
Li1.5Al0.5Ge1.5(PO4)3	NASICON	sputtering	0.25	1.29·10−6	[256]
Li0.5Al0.29Ti1.35P2.9 (LATP)	NASICON	sputtering		6.5·10−6	[257]
Li2O-SiO2	LISICON	PLD		4.1·10−7	[258]
Li3.25Ge0.25P0.75S4	thio-LISICON	PLD	0.38	1.7·10−4	[259]
Li7La3Zr2O12 (LLZO)	Garnet	CVD	0.50	4.2·10−6	[260]
Al-doped LLZO	Garnet	ALD	0.63	1.0·10−8	[261]
Li6BaLa2Ta2O12	Garnet	PLD	0.42	2.0·10−6	[262]
Li2O-V2O5-SiO2(LVSO)		PLD	0.54	2.5·10−7	[263]
LiAlF4		ALD		3.5·10−8	[264]
LiAl0.3S		ALD	0.48	2.5·10−7	[265]
Li2S-P2O5		PLD	0.45	7.9·10−5	[266]
Li3PS4		PLD	0.47	5.3·10−4	[267]

.

3.4.2. NASICON

The general formula of an NASICON-type solid electrolyte is MA₂ (BO₄)₃ (M=Li, Na, K or Ag; A=Ti, Zr, Ge or V; B=P, Si or Mo) (Figure 7b) [268]. In 1976, Goodenough et al. [269] obtained NASICON-type Na_1+xZr₂P_3−xSi_xO₁₂ by a solid-state reaction method at a high-temperature. It comprises the common vertex of the PO₄ tetrahedron and ZrO₆ octahedron [270]. Li-ion solid electrolytes with a NASICON structure can be obtained by replacing Na with Li. Compared with other M ions, LiTi₂(PO₄)₃ and LiGe₂(PO₄) have a higher ionic conductivity, which is, however, still very low. In order to further increase the ionic conductivity, Ti is partially substituted. Li_1+xAl_xTi_2−x (PO₄)₃ (LATP) and Li_1+xAl_xGe_2−x (PO₄)₃ (LAGP) were obtained by partially substituting Ti and Ge with Al. This can achieve a high conductivity of 10⁻⁴–10⁻³ S·cm⁻¹ at room temperature [271]. The investigation of RF magnetron sputtering has been carried out to achieve high conductivity in the thin film form for these materials. The amorphous films possess an ion conductivity of approximately 3·10⁻⁶ S·cm⁻¹, which is equivalent to the LIPON film.

However, the electronic conductivity is noteworthy as it measures at about 10⁻¹⁴ S·cm⁻¹. Elevating the substrate temperature to 300 °C during the sputtering deposition would significantly enhance the conductivity to 2.5·10⁻⁵ S·cm⁻¹ at 25 °C, which is ascribed to the formation of more compact films. Additionally, introducing nitrogen through a N₂ plasma is another approach for enhancing the conductivity of films derived from NASICON targets, similar to LiPON deposition. The resulting amorphous films have reported room temperature conductivities of 10⁻⁵ S·cm⁻¹ [272]. Other amorphous LiZr₂(PO₄)₃ films with a NASICON structure have been deposited by sputtering [273]. The deposited films had a large ohmic resistance and behaved similarly to an insulator. However, the film underwent partial crystallization and transformed into small particles after annealing above 800 °C, leading to a reduced room-temperature ionic conductivity of 1.3·10⁻⁷ S·cm⁻¹. In addition, the chemical compatibility of NASICON materials with metallic Li is a crucial challenge. The Li-metal is highly reactive, which can react with Ti⁴⁺ to form highly resistive interphases [274]. It has been reported that bulk LiZr₂ (PO₄)₃ is unstable when in contact with Li-metal, resulting in a low-resistance interlayer [275]. Therefore, addressing the (electro)chemical compatibility of NASICON electrolyte films in contact with Li-metal is urgent.

3.4.3. LISICON and Thio-LISICON

Bruce and West first described LISICON Li₁₄ZnGe₄O₁₆ and related systems (Li_2+2x Zn_1−xGeO₄) in 1983 [276,277]. The LISICON framework is related to the γ-Li₃PO₄ crystal structure (Figure 7c), which exhibits a relatively low ionic conductivity. Even in doped compounds, such as γ-Li_14.4V_1.6Ge_2.4O₁₆, the room-temperature ionic conductivity remains relatively low, at about 10⁻⁶ S·cm⁻¹. Furthermore, LISICON Li₁₄ZnGe₄O₁₆ is highly reactive with metallic Li and atmospheric CO₂, and the conductivity decreases with time.

It was found that the oxygen ions are replaced by larger and more easily polarized sulfides in the framework, which improves the mobility of Li⁺. Kanno et al. first introduced the Thio-LISICON family to enhance the Li⁺ conductivity [278,279]. Many Thio-LISICON structured solid electrolytes with high conductivity are obtained by aliovalent substitution in solid solutions. Among them, the highest conductivity is apparently observed with Li_3.25Ge_0.25P_0.75S₄. Ohta et al. used PLD to deposit the Thio-LISICON-structured Li_3.25Ge_0.25P_0.75S₄ electrolytes (Figure 7d), which demonstrated an incredibly high room temperature conductivity of 10⁻⁴ S·cm⁻¹ [259], which is roughly an order of magnitude lower than the bulk conductivity of the material. Such behavior was ascribed to the enrichment of Li-content in the deposited films, which leads to a reduction in the number of vacancies available to increase the mobility of Li-ions. However, this material is unstable in bulk form. Later, Ge was replaced by Si to improve its stability and ionic conductivity. The new Thio-LISICON Li_3.4Si_0.4P_0.6S₄ exhibits an even higher ionic conductivity of 6.4·10⁻⁴ S·cm⁻¹ at 25 °C [279].

3.4.4. Garnets

The chemical formula of the garnet structure is A₃B₂ (XO₄)₃ (A=Ca, Mg, Fe, Mn; B=Al, Fe, Cr, Ti, Zr, V), where A, B, and X have 8, 6, and 4 oxygen coordination (Figure 7e). In 2003, Thangadurai et al. [280] first produced Li₅La₃M₂O₁₂ (M=Ta or Nb), exhibiting a conductivity of only 10⁻⁵ S·cm⁻¹ and an activation energy of 0.6 eV. In 2007, Weppner’s group obtained the cubic phase Li₇La₃Zr₂O₁₂ (LLZO) through the doping substitution of the M-site elements, which exhibits high Li-ion conductivity 3·10⁻⁴ S·cm⁻¹ and compatibility with Li-metal [281]. Therefore, LLZO has become a highly promising solid-state electrolyte, which has led to extensive research for high-energy LIBs [106]. LLZO can crystallize into two different types of crystals: a low-temperature tetragonal symmetry (∼10⁻⁶ S·cm⁻¹) and a high-temperature cubic symmetry (∼10⁻⁴ S·cm⁻¹). Numerous techniques have been used to stabilize the high-temperature cubic phase because of the high Li⁺ conductivity. The aliovalent substitutions at the cation site are an illustrative example [239,282]. In the last decade, LLZO has been synthesized with different cationic substitutions via vacuum deposition approaches.

For example, LLZO films were fabricated by pulsed laser annealing and high process temperatures without elemental substitutions. The obtained LLZO films exhibited a conductivity of 1.61·10⁻⁶ S·cm⁻¹ at 25 °C [283]. The Al-doped LLZO films have also been deposited by PLD to study the impact of the Al³⁺ concentration on the conductivity of the material in this paper. The LLZO film with (111) orientation and a moderate level of Al³⁺ incorporation (less than 0.6 moles replacement for Li⁺) exhibited a room temperature conductivity close to 1·10⁻⁵ S·cm⁻¹. In addition, Saccoccio et al. deposited LLZO films doped with Ta element on MgO (100) substrates [284]. The cubic phase films can be synthesized at 50 °C. However, the deposition rate is slow, despite the enhanced conductivity achieved at high laser fluence, extrapolated to 4.4·10⁻⁶ S·cm⁻¹ at 25 °C. In addition to PLD, many other deposition technologies also have been used to synthesize LLZO thin films, such as CVD, ALD, and PVD. Kalita et al. synthesized an amorphous LLZO film with an exceptional conductivity of 4·10⁻⁷ S·cm⁻¹, which is very smooth, dense, and homogeneous [285]. After being doped with Al and Ta, the LLZO films exhibit an impressive conductivity of 1.2·10⁻⁴ S·cm⁻¹ and an activation energy of 0.47 eV, making them a promising candidate for LIBs [286]. Garnet electrolytes have emerged as the most extensively investigated and promising inorganic SSEs for the next-generation commercial TFLIBs.

3.4.5. Sulfides

Sulfide-based materials possess the highest conductivity among all solid electrolyte materials, such as Li₁₀GeP₂S₁₂ and Li_9.54Si_1.74P_1.44S_11.7Cl_0.3, with ionic conductivities of 1.2·10⁻² and 2.5·10⁻² S·cm⁻¹ at 25 °C, respectively, which are comparable to the traditional alkyl carbonate liquid electrolytes in LIBs [11,287]. These high conductivities result from the weak bonds between the less electronegative sulfur anions and Li⁺ coupled with a more accessible microstructural framework [288]. Recently, the sulfide-based electrolytes Li_3+2xPS_4+x have been successfully applied in TFLIBs. Li_3.42PS_4.21 was designed by introducing an excess of Li₂S in Li₃PS₄ by PLD, which showed the most impressive ionic conductivity of 5·10⁻⁴ S·cm⁻¹. They were combined with LCO and Li to establish the LiCoO₂||Li_3.42PS_4.21/Li₃PO₄||Li full TFLIBs. These batteries were able to cycle reversibly at 10 C. The sulfide electrolyte has been demonstrated to be fairly stagnant when in contact with Li-metal because Li₃PS₄ can react with Li-metal to synthesize the electronically insulating Li₃P and Li₂S interphase layers [289]. However, the main issue with sulfide electrolytes is their high reactivity with air and moisture. The entirety of the examined sulfide electrolytes that have been referenced required the operation in glovebox-integrated vacuum deposition equipment. This requirement brings more challenges to the ultimate scale-up and the processing of these electrolytes.

4. Architecture Designs

In addition to the performance of the component materials, the structure of the batteries also greatly impacts the performance. The construction structure of batteries determines the efficient surface contact of the active materials and the diffusion length of Li-ions. That highly influences storage capacity and Li⁺-ion kinetics. The optimization of architecture is an essential condition for designing high-performance TFLIBs.

4.1. Planar TFLIBs

A schematic cross-section of typical thin-film batteries described by Bates is shown in Figure 8a [290]. A wide variety of materials can be used as the substrate if they are stable during the subsequent thin-film deposition and heat treatment. They also should have a relatively smooth surface. Each component of the thin-film batteries, current collector, cathode, anode, and electrolyte is deposited from the vapor phase. A final protective film is needed to prevent the Li-metal from reacting with air when the batteries are exposed to the environment. The typical energy densities that can be achieved for these thin-film cells are 3.6 J·cm⁻² (1 mWh·cm⁻²). With the demand for high-energy devices, it is necessary to expand the deposited film area by folding or stacking the batteries into an optimized volume [291]. However, before Li-ion and Li-free cells can be considered for applications, other problems must be solved. As Li-metal is an anode material, the entire battery must be insulated from oxygen and water vapor during testing and use. Consequently, the packaging requirements are high.

Due to the structural limitations requiring complex mask systems to prepare functional layers of different shapes, Nakazawa et al. designed a vertical FLIB structure by DC sputtering and radio frequency sputtering methods. Li-free batteries are a special type of Li-ion battery that have recently been demonstrated by Neudecker. In such batteries, the Li-anode is formed in situ during activation. This is achieved by electroplating a Li-film at the interface between the current collector (e.g., Cu) and electrolyte (LiPON) [292]. These batteries exhibit high discharge rates and reversible capacities similar to the LIBs with a cLiCoO₂ cathode. In addition, the fabrication, handling, and storage of these batteries are greatly simplified. They do not degrade due to solder reflow, a typical electronic mounting technique, including a short heating cycle to 250 °C. However, due to the small distance between the electrodes and SSEs, the probability of short circuits in the vertical battery structure is relatively high. Therefore, the current research commonly used is the classic cross-section thin film structure to prepare planar TFLIBs. In order to solve the problems of the cross-section structure thin film batteries and improve the energy density, many high-performance anode films and other component film materials have been used to replace Li-metal.

Figure 8. Scheme of different batteries architectures. (a) Planar TFLIBs; (b) Bipolar LIBs [293]; (c) Anode-free TFLIBs; (d) 3D TFLIBs [112]; (e) Flexible and (f) stretchable LIBs [294].

Figure 8. Scheme of different batteries architectures. (a) Planar TFLIBs; (b) Bipolar LIBs [293]; (c) Anode-free TFLIBs; (d) 3D TFLIBs [112]; (e) Flexible and (f) stretchable LIBs [294].

4.2. Bipolar TFLIBs

A bipolar architecture of ASSBs can offer substantial benefits in the battery performance and cost compared to conventional LIBs. The bipolar design refers to the battery configuration, in which unit cells are in direct contact with each other and are connected in series in a batteries module without using external electrical connections (Figure 8b) [293]. The bipolar design allows a simple series connection of multiple batteries to establish the total output voltage and shorten the electronic path in the module, which leads to an increase in power capability [295]. Moreover, the bipolar design can avoid a number of the passive components and parts required for packaging; The external electrical connections are also absent. Without these components, the bipolar ASSBs design can reach the energy density of large-scale battery systems and save costs. The idea of bipolar battery stacking originates from the voltaic pile at the advent of electrochemical science. In 1897, Tribelhorn patented the first bipolar lead-acid batteries, in which a stack of cup-shaped electrodes was separated by glass balls [296]. Kapitza then constructed a bipolar Pb-acid battery in 1923 by immersing Pb plates in an H₂SO₄ solution and achieved a maximum power density of 35 kW·L^–1 and 6 kW·kg^–1 [297]. However, the advancement of bipolar lead-acid batteries with liquid electrolytes has been hindered by short-circuiting problems until advanced electrolyte technology was designed to fix acids in gel or absorbent glass mats.

Inspired by this, bipolar solid-state LIBs have attracted broad interest in recent years. They can improve the output voltage and power density of the batteries. The SEs can eliminate the internal short-circuiting problem of bipolar batteries caused by electrolyte leakage. Shen and Halpert studied the potential advantages of bipolar cell structures from electrode thickness, porosity, bipolar plate thickness, and conductivity [298]. Their calculations indicated that bipolar batteries could exhibit ~40% higher energy and power densities than monopolar batteries. This was mainly due to the volume and weight saving of the passive materials. However, the development of bipolar battery structures still faces some technical challenges, such as the possibility of internal short-circuits between cells, the corrosion sensitivity of the bipolar plates, and the complex manufacturing process for accurate and reliable cell stacking. In order to realize bipolar LIBs, a series of effective strategies are proposed. First, developing SEs with high Li⁺ conductivity and wide electrochemical windows is very important; then, a stable conformal interface (i.e., by interface engineering) between SEs and active electrodes must be built [299]. In addition to developing highly conductive selenium, the biggest challenge is to form a stable conformal Li⁺ interface to promote the transport between the selenium and active electrode materials. Moreover, a bipolar SSLB construction requires electrically conductive bipolar plates (current collectors) that work as an electrical interconnection and partition between the adjacent cells [298]. The bipolar plates may be electrically insulated around their edges to prevent internal short-circuiting between cells.

Although these strategies can significantly enhance the electrochemical performance of the cells, stacking multiple cells in a single package is very challenging. There are only a few reports on bipolar LIBs. The stacking strategies of bipolar LIBs reported in the literature can be roughly divided into the lamination of free-standing SE sheets and the printing of SE slurries. For example, Gambe et al. reported bipolar LIBs via lamination of the LiFePO₄ composition and Li-anode with a quasi-SE composed of a Li–glyme complex and oxide nanoparticles (SiO₂, CeO₂, Al₂O₃, and ZrO₂) [300]. Coin-type bipolar LIBs were constructed by stacking the cells with SUS bipolar plates (Figure 9a). The two- and three-cell stacks exhibited plateau voltages of 6.8 and 10.2 V, respectively. They delivered capacities as high as ~156 mAh·g^–1 at 35 °C and 0.1 C, implying a positive electrode utilization of ~91% (Figure 9b). Furthermore, bipolar LIBs showed high CE (~99%) and excellent cycling stability (91% after 200 cycles). This has primarily been attributed to the forming of Li⁺- conductive interfaces between LiFePO₄ and quasi-SE in the composite positive electrode.

The printing processes of slurries (inks) are simple, low-cost, and easily scalable. Moreover, the printable slurry readily adapts to the shape of the substrate. Then, a novel battery architecture can be implemented with various shapes, shape factors, and functions. This is suitable for the manufacturing of bipolar architectures [290]. Yoshima et al. [301] proved that 12 V-class bipolar LIBs assembled with a thin hybrid SE composed of LLZO and PAN-based gel polymer (GP). LiMn_0.8Fe_0.2PO₄–PAN-GP (Cathode) and Li₄Ti₅O₁₂–PAN-GP (Anode) layers were formed via slurry coating on opposite sides of an Al bipolar plate. A thin hybrid SE layer of ~3 μm thickness was then fabricated by spray-coating a slurry of LLZO (94 wt.%), PVDF binder (2 wt.%), and PAN-GP (4 wt.%) on the cathode and anode (Figure 9c). The Li⁺ conductivity of the hybrid SE is ~6.0·10^–3 S·cm^–1 at 25 °C, which is near to the conductivity of liquid electrolytes. The results show that the coulomb efficiency of bipolar LIBs is over 99.8%, and a stable cycle performance of 85% over 200 cycles is obtained. It also exhibited capacity retentions of ~100% and ~50% at 80 °C and 40 °C, respectively, with respect to the nominal capacity at 25 °C. This demonstrated a successful operation in a wide temperature range (Figure 9d). The bipolar architecture of TFLIBs is a promising solution to improve both the voltage and energy density, thus leading to the development of high-energy batteries in the next generation. However, the development activities in bipolar TFLIBs are still in their early stage. Many technical challenges, such as the stacking design and its manufacturing process, remain to be improved.

4.3. Anode-Free TFLIBs

Li metals can easily react with electrolytes due to their high reactivity. This implies electrolyte consumption, continuous corrosion, uncontrollable dendrite growth, and the formation of irreversible SEI passivation layers [302]. These issues seriously affect the safety, coulombic efficiency, and deteriorate cycling life of batteries. Furthermore, the excessive Li-metal used in the anode decreases the actual specific and volumetric capacity. Inspired by the Li-free batteries structure proposed by Neudecker, anode-free or anode-less rechargeable batteries (Figure 8c) that eliminate the use of excess Li-metal have gained significant attention in the field of LIBs due to their high energy and power density [303,304]. In addition, they bring significantly reduced costs and simplified manufacturing procedures due to the absence of anodes. For example, Zhang’s group further designed anode-free rechargeable Li-metal batteries (Al||LiFePO₄||Cu). They highlighted the promising prospects for their practical applications [305]. Inspired by their fascinating merits, especially their high-energy density, anode-free batteries have sparked the scientific community’s interest.

Dahn’s group [303] recently developed dual-salt LiDFOB/LiBF₄ electrolytes for high-performance anode-free Li-metal pouch cells. Zhang et al. [304] published news and views in Nat. Energy clarifying the fascinating prospects and significant research value of anode-less batteries. Despite numerous advantages, a limited Li source in anode-free batteries is a huge challenge. Woo et al. [306] assembled a LiNi_1/3Mn_1/3Co_1/3O₂||Cu anode-free cell. However, in the initial cycle, only 23% of the extracted Li-ion recovered back into the Li_1−xNi_1/3Mn_1/3Co_1/3O₂ cathode. The sharp capacity loss is mainly attributed to the successive formation of SEI film, “dead” Li-metal, and high impedance interfacial layer. All these processes exhaust the limited Li sources [307]. Therefore, the coulombic efficiency is a crucial challenge for anode-free Li-metal batteries. Recently, several strategies have been developed to enhance the coulombic efficiency and to extend the lifetime of anode-free batteries; this also includes optimizing the current collector and electrolytes, designing favorable protocols, etc.

Assegie et al. [308] successfully grew ultrathin graphene artificial layers on Cu through a more common CVD method. The multilayer-graphene (MLG) film on the Cu is an artificial SEI layer. It stabilizes the electrode interface, encapsulates the deposited Li-metal, and effectively suppresses Li-dendrite formation. This is attributed to MLG films effectively relieving the infinite volume change induced upon cycling. Based on this new design, the anode-free cell with MLG and LiFePO₄ cathodes (areal mass loading: 12 mg·cm⁻²) can deliver a capacity retention of 61% after 100 round-trip cycles, which exhibited a 16% higher retention than the full batteries housing a bare Cu electrode. In addition to constructing an artificial SEI layer or function layer on CCs, pursuing alternative current collectors, such as Li-alloyed CCs, to replace Cu or Al completely and designing 3D structure CCs are also effective methods to improve the coulombic efficiency.

Developing outstanding functional electrolytes is also essential for advancing anode-free LIBs. The SSE is still plagued by premature short-circuiting and huge interfacial resistance attributed to poor contact between the Li-metal and electrolytes. As discussed in the previous section, introducing the interlayer is the most common method to improve the electrochemical performance. Kim et al. [309] explored the impact of the geometry, interlayer, and interlayer species on the LLZO solid electrolyte surface on the plating/striping behavior and performance of anode-free batteries. By contrast, LLZO with a thin Au layer formed a Li–Au alloy layer, inducing fast diffusion. Lithiophilic interlayers on solid electrolytes can effectively induce uniform electric fields and Li-ion flux.

Moreover, the interlayer species (i.e., Au, Ag, and Si) influence the Li-ion distribution and subsequent deposition morphology due to the kinetics of the alloying and precipitation reactions. The interlayer would decrease the interfacial energies by developing a chemically compatible interface. This interface can serve as a buffer layer for Li-ion redistribution and a seed layer for Li-precipitation. Hence, to induce a uniform deposition and alleviate the dendrite problems, solid electrolytes with a uniform morphology and lithiophilic interlayer are needed.

4.4. 3D TFLIBs

Developing new materials and architectures offers an approach to achieving higher energy density in planar TFLIBs. However, it still cannot provide enough power for various small-size applications due to geometric limitations. For conventional 2D planar TFLIBs, it is very complex to combine high power with a high energy performance due to the limited footprint area and volume in cross-section planar TFLIBs. Two routes can be followed to obtain high storage capacities. One is to enlarge the footprint area of 2D all-solid-state LIBs. However, this is not compatible with the small size of integrated devices. Another way is to increase the thickness of the battery’s electrodes, which will significantly increase the diffusion path of Li⁺, limiting the power [310]. Considering these limitations, new architectures are proposed, which are all based on exploring a three-dimensional (3D) geometry (Figure 8d) [112].

One of the advantages of 3D batteries is that the internal surface area between the cathode, electrolyte, and anode is enlarged. This means that much higher currents can be applied, with almost the same amount of packaging and substrate material. Thus, higher power is achieved. Compared with planar solid-state thin film batteries, the 3D TFLIBs improve the volumetric storage capacity, energy density, and rate capability. In addition, the 3D TFLIB architecture introduces a vertical structure into one or both of the anode and cathode. The vertical structures provide electronic bridges, minimizing the ionic migration distance. This enables operation with high power densities, while the volumes of electrode material are still large, and the energy density is high. Introducing 3D structures also reduces the local current density at a given C rate; this improves the performance by reducing the overpotentials and improving the charge transport. In recent years, many 3D structures have been designed for 3D LIBs to improve the energy density. This section will discuss the manufacturing technologies of novel 3D structures and the recent progress in 3D TFLIBs.

The electrochemical performance of microbatteries is highly dependent on the inherent properties of the electrochemically active cathode and anode materials. Compared to conventional planar batteries, the homogeneity of the battery layers is the most critical problem in the three-dimensional case. The poor adhesion between the electrodes and electrolyte layers caused by the fabrication and assembling of the batteries is a serious obstacle [311]. Hence, ALD has become a popular choice for fabricating 3D thin-film electrodes due to the benefit of having high-quality and conformal thin-film electrodes. The following section will review the recent progress of novel 3D structured electrodes, solid electrolytes, and full thin-film batteries.

4.4.1. 3D Cathodes

The previous section introduced the excellent cathode materials for planar TFLIBs, such as layered LiCoO₂, high-voltage cathode materials LMn₂O₄, and high-capacity NCM cathode materials. The cathode materials applied in 3D TFLIBs are similar to conventional cathode materials. The active cathode materials have been applied in 3D TFLIBs, including LiFePO₄, LiCoO₂, LiMn₂O₄, LiNi_xCo_yMn_1−x−yO₂ layered oxides, V₂O₅, MoS, and others. Among them, LiCoO₂ is the most widely used. Shaijumon et al. [312] first designed vertical arrays of Al nanorods directly grown onto Al foil using the pulsed-potential electrodeposition technique. A 30 nm LiCoO₂ thin film was deposited through the thermal decomposition of sol-gel precursors evenly spray-coated onto Al nanorods (Figure 10a). The 3D LiCoO₂ electrode exhibits a high capacity (120 µAh·cm^-2 at C/5), excellent capacity retention, and negligible structural deformation during cycling. The 3D matrix of electrodes meets both the requirements of short transport lengths and large capacity. Later, an even higher footprint capacity for a Li-free 3D electrode was achieved using amorphous VO₂ deposited by ALD on a Pt/TiN-coated Si micropillar array (Figure 10b). A capacity of 130 µAh·cm⁻² at a rate of 0.7 C was observed for such a configuration [313]. Remarkably, rate-performance experiments have indicated that around 40% of the capacity was accessible with a current density of 40 C (50 µAh·cm⁻²). In addition, Labyedh et al. [314] recently prepared LMO thin films on Ni-coated silicon pillar arrays using the solid-state conversion technique (Figure 10c). The LMO thin films on 3D electrodes reached a volumetric capacity of 1200 Ah·L⁻¹, which is close to the theoretical capacity of spinel LiMn₂O₄. It also showed an impressive rate performance in contrast to the planar film. In addition, a 21 times higher capacitance was achieved with the 3D LMO thin film than with the planar one due to the area enhancement offered by the microstructured substrate. In addition to these periodic 3D structured cathode thin film materials, aperiodic 3D structures have also been researched. Shui et al. [315] conducted a synthesis and comparison of sponge-like porous, fractal-like porous, and compact LiMn₂O₄ thin films on Pt foils. The results show that the mesoporous spongy film exhibits the best performance due it having the highest surface area. Wang et al. [191] presented their discovery of 3D porous V₂O₅ nanoparticles with a “multideck cage” structure. The specific capacity of the cathode is 142 mAh·g⁻¹ at 0.5 C and 86.7 mAh·g⁻¹ at 56 C within a voltage window of 2.5 to 4.0 V (corresponding to one Li-intercalation).

4.4.2. 3D Anodes

The active anode materials include graphite, Si, tin and its alloys, and LTO. The most widely reported 3D thin-film anode is TiO₂. The highest footprint capacity for a 3D TiO₂ thin film was achieved by Létiche et al. [317], who combined a Si-microtube array with 155 nm of TiO₂ and covered it with a Li₃PO₄ SSE film. They obtained 370 µAh·cm⁻² at a rate of C/16. Unfortunately, only about 20% was accessible at an applicable rate of 2 C. A slightly lower footprint capacity was achieved using chlorine-doped TiO₂ deposited on Si pillars (242 µAh·cm⁻²). In this case, about 80% was accessible at 2 C; only 37% was accessible at a current rate of 20 C [318]. The higher rate performance is ascribed to the inherently faster Li-ion kinetics due to chlorine doping. This shows the importance of obtaining a large specific capacity with an enlarged surface area and using high-rate performance electrode materials. Therefore, silicon and tin are also widely used in 3D thin film batteries.

As early as 2011, a honeycomb 3D silicon anode material was designed by Notten’s group [119]. The Si honeycombs can retain a stable capacity close to the expected value. The result showed that a suitable 3D structure (Figure 5c) could increase the capacity and improve the Si anode’s stability during cycling. Apparently, such structures can adapt to the stress caused by volume expansion. By optimizing the 3D structure of the Si anode, the high capacity of Si can be maintained, and the stability can be improved. This provides an effective strategy for the practical application of a Si thin film anode. Kim et al. proposed a 3D Si electrode structure formed on check-patterned Cu current collectors with microscale periods fabricated using wire electrical discharge machining (Figure 10d) [316]. The findings indicated that the capacity and cycling performance of the Si electrodes with patterns were relatively high, which could be attributed to the stress relief of the Si anode. In addition, various high-capacity conversion anodes have been deposited on 3D current collectors via ALD, such as RuO₂ [148], MoO_x [147], Fe₂O₃ [265], and ZnO [319]. Dhara et al. developed 3D CNTs@MoOx composites, enabling a much higher areal capacity, about 27 times higher than planar MoO_x. They also found that the thinner MoO_x shows better cycling stability. The optimized MoO_x (300 cycles, ~25 nm)/CNTs could display a stable reversible capacity of 915 mAh·g⁻¹ (areal capacity of 645 µAh·cm⁻²) with an excellent capacity retention after 50 cycles (97.5% compared to the second cycle) [147].

4.4.3. 3D Electrolytes

Solid electrolytes are one of the most crucial components in 3D TFBs. Electrolyte layers provide electrical insulation between the positive and negative electrodes while promoting Li-ion transportation. The material conformality and uniform thickness are critical for uniform resistance over the whole stack. Moreover, a single pinhole will short-circuit the cell, making it useless. Therefore, a suitable thickness that guarantees its proper functioning must be chosen. On the other hand, the electrolyte layer does not contribute to the cell energy density. Therefore, any reduction in the thickness will thus result in an increased energy density. A self-limiting technique is required to achieve good conformality; the most common method used for this is ALD. The conformality originates from self-limiting surface reactions that are separated in time. Many SSE materials deposited by ALD and their most relevant properties have been reported. Among them, the most commonly utilized solid electrolyte for 3D TFLIBs is LiPON. It has a wide electrochemical stability window (0 to 5 V vs. Li⁺/Li), low electronic leakage (<10⁻¹⁴ S·cm⁻¹), and good ionic conductivity (10⁻⁶ S·cm⁻¹). It can be scaled down to nanometer thicknesses while maintaining good insulating properties [320].

In 2015, Nisula et al. [321] fabricated high-quality LiPON films on a 3D-microstructured silicon substrate by thermal ALD. The ionic conductivity of the LiPON electrolyte deposited at 330 °C was measured to be 6.6·10⁻⁷ S·cm⁻¹ at room temperature. In addition, LiPON was also applied to full 3D thin film batteries, indicating it is a promising electrolyte for 3D TFLIBs. However, there are still some problems that hinder its application, including its reactivity with air, which results in the creation of Li₂CO₃ and increased resistance. Another problem is its temperature stability. LiPON is only a good conductor in the amorphous state and cannot withstand high temperatures. Therefore, other high-performance solid electrolytes have been investigated in 3D TFLIBs, such as LLZO and LLTO. Kazyak et al. [261] deposited LLZO on a Si trench substrate by ALD. The resulting LLZO layer had a uniform thickness throughout the entire depth of the trenches, with a step coverage of over 95%. This provided a pathway toward 3D TFLIBs based on the superionic garnet SSEs. As deposition technologies continue to evolve, numerous high-performance solid-state electrolytes can be more effectively integrated into 3D TFLIBs, further enhancing the performance.

4.4.4. 3D Thin-Film Batteries

Despite the significant work exploring the growth and electrochemical performance of individual battery components for 3D batteries, there are few published reports of full 3D solid-state thin film batteries. Various active materials have different chemical properties, and the deposition conditions are also different. Depositing another substance on one active material may cause a chemical reaction at the interface with the active material, thereby affecting the battery performance. In addition, it is also a severe problem to ensure the conformality of each active material, especially for solid-state electrolytes, which may short-circuit the batteries. To date, only a few 3D all-TFLIBs have been reported. Létiche et al. designed the first 3D lithium-ion all-TFLIBs [317]. They created a 3D TiO₂ thin-film electrode with a Li₃PO₄ SSE film coated on the surface. The structure and performance of this 3D thin-film electrode are shown in Figure 10e. The 3D substrate is a double-microtube structure made of Si wafers using a standard photolithography process. An Al₂O₃ barrier was deposited on Si microtubes. Then, Pt is deposited on top of it as a current collector. TiO₂ and Li₃PO₄ films were deposited by ALD. The electrochemical performance shows that this is one of the highest footprint capacities reported for 3D thin-film batteries, with a good cycle capacity of 370 µAh·cm⁻² at a C/16 rate. Unfortunately, applying a higher C rate will significantly reduce the accessible capacity, with only 20% accessible at 2 C. This has been attributed to the poor rate performance of the anatase TiO₂ and the low ionic conductivity of Li₃PO₄. The following significant development are the full complete conformal 3D solid-state TFLIBs made by Pearse et al., as shown in Figure 10f [192]. It was constructed by five conformal layers grown via ALD. The 3D substrate was a periodic array of microholes etched in a silicon wafer. The 3D TFB consisted of 40 nm Ru (current collector), 70 nm V₂O₅ (cathode), which was electrochemically pre-lithiated to form crystalline LiV₂O₅, Li₂PO₂N (SSE), 10 nm SnN_x (anode), and 25 nm TiN (second current collector). Cu dots were deposited on top of the stack using electron beam deposition. Planar solid-state batteries made from the LiV₂O₅-SnN_x couple exhibit stable capacities of 2.6 μAh·cm⁻² (37 μAh·cm⁻²·μm⁻¹ normalized to the cathode thickness) for hundreds of cycles. By integrating with 3D substrates, the surface discharge capacity of these cells has been expanded to 9.3 times that of planar cells. Most importantly, the 3D structure improves the rate performance and coulombic efficiency while increasing the footprint capacity. These results indicate that the 3D architectures simultaneously increase the energy and power density, which are intrinsic to planar architectures with the trade-off between energy and power. 3D all-solid-state TFLIBs can indeed effectively increase the energy density due to the enlarged surface area. However, there are still significant obstacles to overcome before these batteries can be used in commercial applications. More research and development are necessary to address the remaining challenges. Among them, the choice of composition materials, 3D structures, and deposition technology must be further optimized.

4.5. Other Novel TFLIBs

At present, flexible and stretchable Li-ion batteries (Figure 8e,f) are the ideal power sources for wearable electronics, owing to their high energy density, long cycle life, and excellent rate capability [322]. Flexible batteries have a similar principle to conventional batteries because each contains an anode, cathode, separator, and electrolyte. However, conventional batteries with rigid and fragile planar structures cannot effectively meet the requirements of wearable electronics. Therefore, it is crucial to produce flexible batteries with substantial flexibility while maintaining a high energy density and long-term stability [323]. So far, several approaches have been proposed for developing flexible micro batteries using different configurations, such as a paper-like configuration [324], bridge-island batteries designs [325], sponge/porous configurations [326], and wavy structures [327]. Paper-based TFLIBs use the surface roughness and porous structure of paper and textiles, which are ideal for ion transportation. At the same time, the flexibility of paper and textiles makes it possible for batteries to bend. Cui and co-workers [328] demonstrated the first flexible-paper LIBs, which employed paper as the separator and independent carbon nanotube (CNT) films as the CCs for both the anode and cathode. These microbatteries showed good electrochemical and mechanical performances. It displayed a specific capacity of 122 mAh·g⁻¹. The coulomb efficiency in the first cycle is 85%, which increases to 93% after 20 cycles. It exhibited an excellent mechanical performance after manually bending the device down to 6 mm for 50 times without failure.

Graphene is a 2D carbon nanostructure with covalently bonded honeycomb array materials. It has been investigated as a material to produce LIB flexible electrodes. Wang et al. [329] reported the first graphene paper LIB in 2009. In this work, they prepared graphene aqueous dispersions from commercial graphite. A super-paper material with a smooth surface and glossy two sides was formed through the vacuum filtration of graphene aqueous dispersion. Microbatteries using graphene paper as the cathode and a Li-metal foil as the anode delivered a discharge capacity of 528 mAh·g⁻¹. Later, Zhao et al. [330] designed a soft Si/graphene electrode for LIBs using acid-sonication treatment to create graphene with carbon vacancies. This composite material showed a reversible capacity of 3200 mAh·g⁻¹ at 1 A·g⁻¹ and high stability over 150 cycles. In addition to these examples, numerous other electrode materials for flexible batteries have been developed [331], including transition metal oxides [332] and sulfides [333]. Other flexible configurations have also been designed for LIBs, such as sponge structures [294]. Optimizing the flexible configuration will significantly improve thin-film batteries’ stability and energy density.

5. Outlook

All solid-state thin-film batteries (TFLIBs) have been produced by various deposition techniques. These techniques efficiently avoid microscopic defects at the solid-solid interface and minimize barriers at the junctions. TFLIBs exhibit high stability, a long cycle life, a wide operating temperature range, and a low self-discharge rate. Recently, fruitful progress has been achieved in TFLIBs. Some TFLIBs have been employed in wearable electronics and low-power microdevices. However, the demand for thinner, lighter, space-saving, and flexible-shaped batteries continues to increase. At present, commercialized TFLIBs cannot fulfill these requirements. Consequently, the creation of high-performance TFLIBs is required to be developed. In this review, the research advances made in recent years for developing high-performance TFLIBs have been outlined. Primary attention was devoted to component materials optimization and cell architecture design.

Exploring novel materials and optimizing the present materials are critical to developing high-performance TFLIBs with improved storage capacity and recyclability. For the anode side, the electrode films should have low redox potential with the limited volume change, low interface resistance, and a high-capacity retention rate. Developing multiple-compound systems, including complex alloys, pre-lithiated anodes, and in-situ deposited Li anodes are also promising strategies for future TFLIBs’ anodes. On the cathode side, the electrode films should, in contrast, have high redox potential with stable and high capacity. Enhancing the stability and enlarging the capacity for high-voltage cathode materials is an indispensable direction to enhance the energy density and cycling life of TFLIBs. The electrolyte films should be highly ionically conductive, lithiophilic, and compatible with both the cathode and anode. In addition, the electrolyte films should be stable to operate at a wide range of voltages. The current collectors should be highly electronically conductive and compatible with the electrode materials. More novelties, such as memory and liquid alloys, should be explored to build superior current collectors for TFLIBs.

With the development of various advanced deposition technologies, such as CVD, PVD, PLD, ALD, 3D printing, surface modification, doping, and nano-structuring, are several common methods to improve the performance of these component materials. Although many advanced composite materials have been reported, few commercially available materials are still available. The high cost of manufacturing TFLIBs is an important issue hindering their widespread use in commercial applications. The following techniques are available to lower the cost of manufacturing TFLIBs: (1) Roll-to-Roll Manufacturing: This technique involves the continuous fabrication of TFLIBs on a roll of substrate material, allowing for high-volume and low-cost production; (2) Low-cost substrate materials: TFLIBs can be produced on numerous substrate materials, such as glass, metal foils, and polymer films. Using cost-effective and widely available substrate materials can reduce the overall manufacturing cost of TFLIBs; (3) Improving Manufacturing Processes: Improving deposition efficiency can reduce manufacturing costs. 3D printing technologies enable the fabrication of TFLIBs using low-cost, high-throughput methods; (4) Materials recycling: Recycling the component materials during the fabrication of TFLBs can also reduce the cost. Lowering the manufacturing cost of TFLBs will make them more purchasable, thus facilitating their widespread adoption in a broad range of applications.

The architecture of the batteries will also significantly affect their properties. Designing 3D TFLIBs will increase the areal energy and power densities. Various 3D methodologies have been proposed to increase the batteries’ storage capacity, while keeping the same footprint area. In this review, the research progress on 3D component materials and complete 3D thin film batteries in recent years has been investigated in detail. Although many achievements have been made in 3D TFLIBs, they have not yet achieved commercial success. Many problems remain, such as the conformity of the 3D thin films, optimizing 3D structures, and lowering the manufacturing costs. Future efforts in material development and micro-manufacturing technology are still needed to realize commercial 3D TFLIBs.

Reliability is a critical factor for TFLIBS and is affected by various factors. The mechanical integrity of TFLIBs is one of the main issues. TFLIBs are very vulnerable to mechanical load, leading to cracking or delamination of the layers within the battery and further performance fade. Advanced materials with higher mechanical stability will help suppress crack and delamination formation in the battery layers. High power is also a reliability concern for thin-film lithium batteries. The cell should be able to operate at a high current to achieve a high energy transformation rate. Working with a high current would lead to overheating, increased resistance, and electrode degradation. This results in a reduced capacity, shorter cycle life, and potential safety hazards such as thermal runaways. Therefore, it is crucial to design TFLIBs with the appropriate electrode materials, thickness, and structure to handle high currents. The optimization of the manufacturing process to ensure uniformity and consistency is also important. In addition, it is desirable to improve the thermal management. The cycling process of the battery must be conducted mindfully to avoid exposure to excessive power demands. This can prolong the lifespan of the battery and reduce safety risks.

With the continuous advancements in high energy density and reliability, TFLIBs can be widely applied in life in wearable devices, medical implants, or smart cards. We anticipate that this review will provide useful direction for the creation of high-energy-density and stable TFLIBs, enabling them to be used in an even wider variety of real-world scenarios.