## Abstract

Li metal batteries (LMBs) are one of the most promising high-energy-density batteries. However, solid electrolyte interphase (SEI) and Li dendrite substantially form in LMBs. Due to low ionic conductivity, inhomogeneity, and poor mechanical and electrochemical stability of native SEI, dendritic Li nucleates and grows, which in turn induces the fracture of SEI and promotes the formation of new SEI, causing the loss of active materials and safety issues for LMBs. Understanding the SEI–dendrite interactions could play a critical role in developing LMBs. For instance, modified SEI has been demonstrated to suppress dendrite growth and improve battery performance. In this short review, we discuss the underlying mechanisms of SEI–dendrite interactions and strategies for improving battery cycling performance.

## 1 Introduction

Lithium (Li)-based batteries are widely used as energy storage devices with the merits of light weight and high energy density [1,2]. Among all kinds of Li batteries, Li metal batteries (LMBs) have been recognized as one of the most promising storage systems for the applications in energy, transportation, and industry. Li metal anode, possessing ultrahigh theoretical specific capacity (3860 mAh g−1) and extremely low electrochemical redox potential (−3.04 V versus standard hydrogen electrode), is considered as “Holy Grail” of negative electrodes [35]. Therefore, LMBs have attracted increasing attention, showing great potential in the next-generation rechargeable batteries.

However, LMBs have safety issues because of dendritic Li, and its formation process includes Li nucleation and growth. In the beginning, Li ions transfer to the anode surface and thus obtain electrons, forming the initial nuclei. Then, the subsequent Li ions preferentially deposit on the initial nuclei due to higher electric field strength induced by the protuberance site of initial nuclei. As the amount of deposited Li increases, dendritic or mossy Li is gradually formed on the surface of Li metal electrode, which is defined as “Li dendrite” [68]. The existence of Li dendrite not only exposes fresh Li to electrolytes, resulting in persistent reactions between Li and electrolytes and irreversible capacity depletion, but also increases the cell impedance and leads to low Coulombic efficiency (CE) [9]. In addition, the microstructural dendrite can easily transform into “dead Li” as a result of losing the electric connection with the Li metal anode. The formation of dead Li also causes severe capacity loss and thus influences the cycling performance of the battery during charge/discharge process [10]. Furthermore, Li dendrites could uncontrollably puncture the separator inside the battery, leading to short circuits and even catastrophic battery failure involving fatal safety issues [11]. Therefore, the safety issues produced by Li dendrite are critical factors in determining whether the LMBs have stable cycling performance. Among the roots affecting the formation of Li dendrite, solid electrolyte interphase (SEI), induced from chemical and electrochemical reactions between Li metal and electrolytes, plays a crucial role in influencing Li deposition behaviors. Natively formed SEI tends to exhibit extremely inhomogeneous properties in its structure and components, which brings about nonuniform Li nucleation on the anode surface, consequently resulting in uneven Li deposition [12]. The formation of Li dendrite is inevitable, and SEI is prone to crack under the stress created by dendrite. Furthermore, Li dendrite beneath the SEI reacts with electrolytes instantly, and new SEI is thus generated [13]. Native SEI is mechanically fragile and could not endure the large volume change of Li metal anode, resulting in cracks of SEI. Therefore, native SEI is unable to suppress the growth of Li dendrite due to its poor mechanical properties, and the repeated SEI fracture and regeneration cause unstable SEI and consumptions of active materials, which degrades the cycling performance in LMBs.

In general, natively formed SEI film is heterogeneous. When Li ions transport through the native SEI, nonuniform Li nucleation and deposition occur on the anode surface, which promotes the formation of Li dendrite. Moreover, the mechanical strength of native SEI is not enough to inhibit the Li dendrite growth. Therefore, Li dendrite can break the SEI film and expose itself to the electrolyte, causing many disadvantages that are unfavorable to the LMBs. In order to suppress the growth of Li dendrite, many efforts have been made to modify the SEI film, and great success has been achieved in this field. In this review, we aim to discuss mechanistic insights into SEI–dendrite interactions and strategies related to dendrite suppression by regulating SEI for safe and stable performance of LMBs (Fig. 1).

Fig. 1
Fig. 1
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## 2 Interactions Between Solid Electrolyte Interphase and Lithium Dendrite Growth

An ideal SEI film should have high ionic conductivity, enough thickness, preferable uniformity, excellent mechanical strength, and good chemical stability. However, the spontaneously formed SEI layer has the inhomogeneity of structure and chemical composition. The SEI inhomogeneity significantly affects its intrinsic properties, such as ionic conductivity, mechanical properties, and electrochemical stability, which are the key factors influencing the growth of Li dendrite. In this section, we explore the underlying mechanisms of Li dendrite growth caused by the SEI film.

### 2.1 Ionic Conductivity.

For charging, Li ions experience three steps to reach the Li anode surface and be reduced. First, the solvated Li ion sheds its solvent molecules and approaches the Schottky vacancy located on the upper surface of the SEI. Then, the Li ion diffuses through the solid SEI film. Finally, it reaches the anode surface and is immediately reduced after obtaining an electron from the current collector [14]. Peled et al. [15] found that the resistance to diffusion of Li ions via SEI grain boundaries (GB) is lower at room temperature. Newman et al. [16] established a model to simulate the SEI growth and Li ion transport through the SEI layer at specific diffusion pathways, and they discovered that diffusion through grain boundaries is a vital diffusion form for Li ions. Borodin et al. [17] developed a force field model by molecular dynamics (MD) method to simulate the dynamics of Li ions within the SEI film, which was modeled without vacancy or interstitial. Based on the density functional theory (DFT) calculations, Ramasubramanian et al. [18] explored Li ion diffusion through the grain boundary of polycrystalline SEI containing three major inorganic components: Li2O, LiF, and Li2CO3, as shown in Fig. 2(a). They found that grain boundary possesses numerous open channels and Li ion diffusion in the channels of grain boundary is universally faster than in the neighboring crystalline regions within the grain interiors. In particular, they analyzed several different GB structures and revealed that the LiF/LiF, Li2O/Li2O, and LiF/Li2O GBs are the most stable configurations in all structures. In addition to grain boundary diffusion, Shi et al. [19] built a double-layer SEI model consisting of porous (outer) organic and dense (inner) inorganic layers, based on which they proposed a new ion diffusion mechanism: pore diffusion in the outer layer and interstitial knock-off diffusion in the inner layer (Fig. 2(b)). Besides, Single et al. [20] also concluded that there are many pathways for Li ions to migrate through the SEI layer, including porous diffusion, grain boundaries, interstitials, and vacancies. The transport mechanism of Li ions through SEI could be considered as a combination of migration and diffusion. Native SEI contains various components, and each component has its specific transport mechanism. For porous organics, Li ions tend to diffuse through loose spaces between the compositions. In contrast, Li ions migrate through vacancies or interstitials for the transport of Li ions in dense inorgaincs. For instance, previous studies show that Li ions prefer to migrate through vacancies rather than interstitial knock-off in LiF [21]. And the direct exchange of adjacent Li ions is the mechanism of migration in Li2O [22].

Fig. 2
Fig. 2
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However, the migration of Li ions possesses its specific regulations regardless of the diffusion pathway mentioned above. Wu et al. [23] mentioned that the migration of Li ions need to break the chemical bonds in the previous site, subsequently forming new bonds at the neighboring site. In particular, the first two steps of the migration constitute a kind of energy called energy barrier (migration activation energy Ea) which Li ions must overcome. Moreover, they also found that if the kinetic energy (Ek) of Li ion itself is higher than the migration activation energy (Ek>Ea), the ion is capable of overcoming the energy barrier and moving to the neighboring site successfully. Otherwise, it is fixed in its original position. They considered the ions with larger kinetic energy than energy barriers as charge carriers. Wu et al. gave the formula for calculating the ionic conductivity of the SEI:
$σ=μnz$
where σ is the ionic conductivity, μ is the ionic mobility, n is the number density of charge carriers, and z is the electric charge of each charge carrier. The formula shows that the ionic conductivity is proportional to the number of charge carriers, which in turn depends on the migration activation energy. The transport mechanism of Li ions mainly depends on three parameters: conductivity, diffusivity, and transference number. As mentioned above, the transference number could be considered as 1 for the inorganic components of SEI. Thus, the SEI conductivity is essentially the only unknown parameter and depends on the SEI migration activation energy, which can be calculated by DFT. Therefore, we can conclude that ionic conductivity and migration activation energy are two important parameters for conducting ions and they are closely related to each other.

Table 1 summarizes the ionic conductivities and energy barriers of several common components of SEI. The ionic conductivity of LiF is approximately 10−31 S cm−1 at 25 °C, while the energy barrier is as high as 0.729 eV [24,25], which is adverse to the conduction of Li ions. LEDC and LMC, typical Alkyl Carbonates salts, have ionic conductivities less than 1 × 10−9 S cm−1 at room temperature [26,27]. In comparison, LEMC has ionic conductivity of 6.4 × 10−6 S cm−1 at 25 °C [28]. The ionic conductivity of Li2O on pure artificial SEI is about 10−9 S cm−1 at room temperature, three magnitudes higher than that of sintered Li2O pellet (10−12 S cm−1) [29,30]. It is found that the ionic conductivity of amorphous Li2O can be three magnitudes higher than that of the crystalline Li2O [32], which demonstrates that the ionic conductivity is structurally dependent. The energy barrier of Li2O in artificial SEI is 0.58 eV. At room temperature, the ionic conductivity of Li2CO3 is in the range of 10−11 to 10−8 S cm−1 [31]. Simultaneously, the energy barrier of Li2CO3 varies from 0.227 to 0.491 eV [24]. Overall, the main components of native SEI have large energy barriers, resulting in low ionic conductivities.

Table 1

Ionic conductivity and energy barriers of components of SEI

ComponentsIonic conductivity (S cm−1)Energy barrier（eV）Ref.
LiF10−310.729[24,25]
Li alkyl carbonates<10−90.76[2628]
Li2O10−90.58[29,30]
Li2CO310−11–10−80.227–0.491[24,31]
ComponentsIonic conductivity (S cm−1)Energy barrier（eV）Ref.
LiF10−310.729[24,25]
Li alkyl carbonates<10−90.76[2628]
Li2O10−90.58[29,30]
Li2CO310−11–10−80.227–0.491[24,31]

Although the ionic conductivity of SEI can be obtained, it is not sufficient to understand the mechanism that SEI influences the behaviors of Li electrodeposition. Vu et al. [33] performed phase field simulations of Li electrodeposition. First, the ionic conductivity of the native SEI was computed using DFT calculations: 5 × 10−7 S cm−1, which indicates that the ionic conductivity of SEI is low.

They modeled Li metal anode covered by the SEI film with low ionic conductivity (Fig. 3(a)). Notably, since the surface of pristine lithium metal is uneven at the nanoscale, a tiny tip was simulated on the Li metal surface. In fact, the tip can be considered as the nuclei of Li metal, and subsequent Li ions preferentially deposit on the surface of the nuclei. In their study, the ionic conductivity of native SEI was set in the range of 5 × 10−7 to 5 × 10−5 S cm−1 under certain current densities, and the thickness of native SEI was set in the range of 100–500 nm. Based on the model, Li deposition morphology could be captured. Meanwhile, Vu et al. also gave a reference standard for dendrite growth: growth-rate ratio Δdtipdvalley (Fig. 3(b)). If Δdtipdvalley > 1, it represents dendrite formation; if Δdtipdvalley < 1, it represents uniform deposition. Figure 3(c) depicts that Li tends to deposit at the tip when the growth-rate ratio is greater than 1, resulting in dendrite formation. The main reason for this phenomenon is the low ionic conductivity of SEI film, which leads to the uneven distribution of the internal electric field and concentration field, causing a huge difference in the number of deposited Li atoms between the tip and the valley. Figure 3(d) shows the distribution of potential. The low ionic conductivity of SEI film leads to a large potential gradient in the thickness direction, and the higher potential at the tip makes it easier to attract Li ions compared to the valley. As a result, Li is preferentially deposited at the tip. Vu et al. also provided a relationship of growth-rate ratio Δdtipdvalley with SEI thickness and conductivity. The value of the growth-rate ratio reaches up to 8.8 for the SEI with an ionic conductivity of 5 × 10−7 S cm−1 and a thickness of 500 nm. As the thickness of SEI is reduced, the growth-rate ratio gradually decreases to 7.0 (dSEI = 400 nm) and 4.0 (dSEI = 200 nm), respectively. Therefore, the low ionic conductivity of SEI layer is more likely to cause the growth of dendrite. In addition, Vu et al. gave the criteria for the ionic conductivity of SEI film that would achieve uniform deposition: a high ionic conductivity greater than 5 × 10−5 S cm−1. For example, Fig. 3(f) shows that the high ionic conductivity of SEI (8 × 10−4 S cm−1) contributes to fully flat Li deposition morphology, distinct from the dendrite morphology induced by the low ionic conductivity of SEI layer. Hence, it can be concluded that high ionic conductivity facilitates uniform flat deposition, while low ionic conductivity promotes Li dendrite growth.

Fig. 3
Fig. 3
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Yoon et al. [34] also conducted phase field simulations for the effect of ionic conductivity of 200 nm-thick SEI on the Li deposition morphology. They initially set an ionic conductivity 100 times lower than the electrolyte. Then, they observed an obvious concentration gradient inside the SEI film and Li deposition morphology with sharp tips, as shown in Fig. 4(a). Due to the low ionic conductivity of native SEI, the diffusion rate of Li ions through the SEI membrane is low, resulting in an insufficient supply of ions on the surface of the Li electrode and a significant concentration gradient within the SEI layer. Under such a concentration gradient, tips are formed on the surface of Li metal seed. Any small tip of Li metal results in a high local current density and then attracts more Li ions preferentially deposited on it, which initiates dendrite growth. If the ionic conductivity is increased to 10 times lower than that of the electrolyte, the concentration gradient in the SEI layer is substantially reduced and the surface of Li seed becomes smoother and flatter compared with the previous case (Fig. 4(b)). It further illustrates the importance of the SEI layer with high ionic conductivity for uniform deposition and dendrite-free morphology.

Fig. 4
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### 2.2 Inhomogeneity

#### 2.2.1 Structure Inhomogeneity.

The thickness of SEI film is in the range of several nanometers to several hundred nanometers [35], depending on the operating conditions of the battery. In general, SEI is inhomogeneous in thickness, as shown in Fig. 5 [36]. The total transport time of Li ions through SEI is less in the thin part than in the thick part, and thus, Li preferentially nucleates to form dendrite where the SEI is thin, resulting in uneven nucleation and deposition. Furthermore, the extremely thin part of SEI efficiently conducts electrons, although inorganic components of SEI are insulating, causing Li ions to be reduced to metallic Li before reaching the Li electrode, which accelerates the growth of dendrite. The thinner the SEI film is, the easier it is to conduct electrons. In short, the thickness of SEI layer is generally inhomogeneous, which leads to the rapid transfer and easy acquisition of electrons of Li ions at the thinner sites of SEI, forming dendritic or mossy lithium. It is worth noting that the thinner parts of SEI are easy to fracture under the stress of dendrite if the SEI does not have excellent mechanical strength. Therefore, in terms of the SEI thickness, a uniform and appropriate thickness can facilitate the transport of Li ions through the SEI and inhibit the contact of ions and electrons. This enables an electrochemically stable interface [37].

Fig. 5
Fig. 5
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In order to explore how the inhomogeneous thickness of SEI film affects its interior stress concentration, Shen et al. [36] performed corresponding phase field simulations under different SEI thicknesses. They gave a dimensionless parameter to describe the inhomogeneity of the SEI structure: “defect depth ratio pd,” which was defined as the ratio of the depth of defect and overall thickness of SEI (hd/hSEI). The larger its value is, the more obvious the defect will be. They built the SEI model under three defect depth ratios pd (0.9, 0.5, 0.1) and simulated the influence on the SEI itself caused by defects or inhomogeneity of SEI thickness (Fig. 6).

Fig. 6
Fig. 6
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It can be seen from the simulation results that the Li metal protrusion presses the SEI during the cycle, resulting in the local stress and SEI deformation, and the maximum stress appears in the center of the SEI defect. The stress increases gradually over time, eventually reaching the yield strength. Most importantly, as the defect depth ratio increases, the phenomenon of stress concentration of the SEI is more prominent, and the failure is faster. That is, the thinner the SEI, the more fragile it is, because the thin site in the SEI is prone to stress concentration, where its mechanical properties are usually weak. Thus, dendrite can easily penetrate the thin part of SEI layer and grow vertically. Liu et al. developed a computationally concurrent coupled model to simulate the interactions between Li dendrite and the SEI film [38]. In their model, the thickness in the central area of SEI is much thinner than other regions, and Li protrusion subsequently forms in the central region due to the fast deposition rate, indicating that SEI inhomogeneity could lead to the formation of Li dendrite. Moreover, they also explored the influence of dendrite on the SEI thickness. The thickness of SEI decreases rapidly at the dendritic tip due to its large curvature. The large SEI thickness not only makes Li dendrite easier to break from the substrate surface but also largely decreases Li electrodeposition kinetics. As the difference in the inhomogeneous distribution of SEI thickness between the dendritic tip and flat region becomes larger, more Li ions prefer to deposit at the tip, making SEI undergo repeated fracture and regeneration. To improve the stability of SEI, a uniform structure is indispensable.

#### 2.2.2 Components Inhomogeneity.

SEI, as a protective film on the surface of Li metal anode, has been the focus of attention, and its chemical compositions are complex and irregular. As early as 1970, Dey discovered that the electrolyte inside the battery reacted with Li to form a thin film [39]. Later, Peled named this thin film as the solid electrolyte interphase in 1979 [40]. Subsequently, Peled discovered that the chemical composition of SEI is mainly composed of Li-containing compounds and polymers, including LiF, Li2O, etc. [15]. Therefore, he constructed a mosaic model to describe the microstructure and composition of SEI (Fig. 7(a)). The mosaic model perfectly reflects the inhomogeneity of SEI in chemical composition and lays a good foundation for future research on SEI. Furthermore, Aurbach et al. [41] found that the SEI was a mixture of dense inorganics inside and porous organics outside, and they proposed a bilayer model showing that the SEI is composed of a wide variety of organic and inorganic species (Fig. 7(b)). Therefore, the complex variability of the native SEI confirms the inhomogeneity in terms of chemical composition.

Fig. 7
Fig. 7
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The SEI contains a variety of insoluble substances, which have different properties. For example, organic components in SEI, such as ROLi and ROCOLi, exhibit excellent electrochemical stability. Compared to inorganic components, organics have lower energy barriers and thus are more favorable for ion diffusion [42]. Among the inorganics, Li2CO3 promotes the formation of Li dendrite, while LiF contributes to uniform Li deposition. It is more difficult for LiF to conduct electrons compared to Li2CO3, which helps suppress the growth of dendrites [43]. Therefore, the chemical inhomogeneity will cause nonuniform ion conduction characteristics within the SEI layer, which leads to uneven nucleation and dendrite growth on Li metal anode. To confirm the effect of inhomogeneity on Li deposition, Ozhabes et al. [44] calculated the surface energies and diffusion barriers of common components of the SEI and analyzed the surface diffusion properties of the components. It is worth mentioning that surface diffusion is one of the most important processes of electrodeposition. Moreover, surface energy and diffusion barrier are two important parameters to determine surface properties. Components with high surface energy possess a sufficiently large resistance to suppress dendrite growth. Ion diffusion rate depends on the diffusion barrier. The larger the diffusion rate, the easier it results in uniform deposition. Components with a low diffusion barrier are less likely to form dendrite.

In Fig. 8, Li halides tend to have low diffusion barriers and high surface energies, which favor the formation of dendrite-free morphology. In contrast, Li2CO3 has a higher diffusion barrier and smaller surface energy, both of which are detrimental to uniform Li deposition. Even Li2O seems to be undesirable due to its high diffusion barrier. In general, due to the inhomogeneous composition of native SEI, each component has its unique surface energy and diffusion barrier. Thus, Li diffusion rate on the surface of different components is disparate. This results in nonuniform deposition and subsequent dendrite growth. In addition to surface diffusion, Li diffusion through the interior of the SEI is also an important process for electrodeposition. In this regard, Hao et al. [45] simulated the effect of inhomogeneity within the SEI on Li electrodeposition. They set the diffusion coefficient in the central region of the SEI as Dc, different from that of D in the other parts, which was achieved by changing the diffusion barrier within the SEI film. Therefore, Dc/D represents the inhomogeneity of the SEI. For the three cases of Dc/D = 1, Dc/D = 3.2, Dc/D = 0.3, they observed different Li deposition morphologies.

Fig. 8
Fig. 8
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In Fig. 9(a), uniform deposition can be achieved if the diffusion coefficient of the central region is the same as that of the other regions, which means that uniform SEI results in the dendrite-free morphology. The larger diffusivity Dc in the central region, however, results in faster Li deposition and dendrite growth in the central region (Fig. 9(b)). Moreover, the Li concentration in the central region also has a clear increasing trend attributed to the larger diffusivity Dc (Fig. 9(d)). The trend becomes more obvious as the diffusivity Dc increases. In contrast, a lower diffusivity Dc in the central region causes a slower deposition rate and decreasing concentration gradient (Figs. 9(c) and 9(e)). It can also induce inhomogeneous deposition morphology. Actually, the inhomogeneity of SEI composition is much more complicated than that in the above simulations. SEI is composed of multiple components, and different components have different diffusion barriers or diffusion coefficients mentioned in Table 1, thus further confirming the complexity of its structure. In short, because of the SEI inhomogeneity, the sites where D is large are more favored to occur Li dendrite growth than that where D is small. Hence nonuniform Li-ion transport exists in the SEI and dendrites are prone to grow in the places with higher D. To sum up, the inhomogeneous compositions of SEI not only lead to uneven Li nucleation but also cause Li dendrite growth during Li plating.

Fig. 9
Fig. 9
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### 2.3 Mechanical Properties.

The mechanical properties of SEI are critical to maintaining LMBs' performance. The SEI with excellent mechanical properties can not only prevent the further reactions of Li metal with electrolytes but also suppress the growth of dendrite. Wang et al. [46] proposed that the growth of Li dendrite originated from the compressive stress generated by the bottom of the anode. Figure 10 shows that compressive stress leads to the vertical growth of dendrite during charging. Moreover, the inhomogeneous SEI leads to stress concentration under the Li nuclei, resulting in the rapid growth of dendrite.

Fig. 10
Fig. 10
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In order to explore the effect of vertically grown dendrite on SEI film, Hao et al. [45] simulated the stress distribution induced by dendrite growth using the finite element method. In the simulation, they set the Young’s modulus and Poisson’s ratio of SEI layer to 10 GPa and 0.3, respectively, both of which are important parameters to characterize the mechanical properties of SEI.

The simulation results are shown in Fig. 11. Li dendrite grows in the x-axis direction, which compresses the SEI in the central region to a large extent, generating large compressive stresses within the SEI. Meanwhile, the two sides near the central area need to bear larger tensile stresses to balance the stress in the SEI. Large compressive stress is the main root leading to SEI fracture. In addition to the stress distribution in the x-axis direction, the stresses generated in the y-axis direction of the central region undergo a significant transition: from compressive stresses to tensile stresses. Therefore, Li dendrite growth can produce large stresses regardless of the x-direction or the y-direction, which could break the original SEI layer, thereby forming fast Li-ion transport channels and promoting Li dendrite growth. In addition, Hao et al. also analyzed the effect of SEI inhomogeneity on the shear stress generated by dendrite through the mechano-chemical coupling. They introduced an electrochemical Biot number, and quantified the inhomogeneity of SEI as γ=Dc/D, the ratio of ion diffusivity in the central SEI region to that of adjacent regions.

Fig. 11
Fig. 11
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In Fig. 12, the inhomogeneity of SEI is the main reason for the increase in its internal shear stress. The value of maximum shear stress is also related to the electrochemical Biot number for a specific SEI inhomogeneity. Therefore, the combined effect of electrochemistry and SEI inhomogeneity leads to larger shear stresses accompanied by Li dendrite growth.

Fig. 12
Fig. 12
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According to the above simulation results, stress generation induced by dendrite growth is the primary mechanism of SEI failure. Nevertheless, the phenomenon of SEI failure caused by dendrite can be avoided if the SEI has excellent mechanical stability. Yu et al. [47] found that excellent mechanical stability included “strong” and “soft.” “Strong” means high shear modulus, which can suppress the growth of Li dendrite and reduce potential safety issues. “Soft” means high elastic modulus, which can ensure uniform and flat Li deposition. The native SEI layer is composed of various components with different mechanical moduli. For example, among the inorganic components of SEI, Li2CO3 has higher Young’s modulus than LiF, which suggests that Li2CO3 has better mechanical stability [48]. Furthermore, Table 2 summarizes the Young’s modulus of several common components of SEI. It can be seen that the Young’s modulus of inorganic components is several orders of magnitude larger than that of the organic components, implying that the dense inorganic components have higher mechanical strength.

Table 2

Young’s modulus of SEI components

SEI componentsYoung’s modulus (GPa)Ref.
ROCO2Li<1[13,49]
LiF65.0[13,49]
Li2CO375.0[13,49]
Li2S82.6[13,49]
Li2O169.0[13,49]
SEI componentsYoung’s modulus (GPa)Ref.
ROCO2Li<1[13,49]
LiF65.0[13,49]
Li2CO375.0[13,49]
Li2S82.6[13,49]
Li2O169.0[13,49]

Although the modulus of inorganic components of SEI is large enough to suppress the growth of dendrite, the overall mechanical strength of SEI is too weak to withstand the huge volume change of Li metal due to the inhomogeneity of SEI. When the stress generated by Li dendrite is greater than the one that the native SEI can endure, the SEI fails and breaks, resulting in irreversible consumption of fresh Li exposed to the electrolytes. Therefore, in addition to ensuring the uniformity of SEI, improving mechanical properties is also a key aspect to suppress dendrite. Theoretical studies show that robust SEI with a modulus larger than 1 GPa can effectively suppress the growth of dendrite [50,51]. In practice, native SEI tends to possess a low effective modulus.

Measuring the modulus of SEI has been widely conducted to explore its mechanical properties. Gao et al. [52] used atomic force microscopy (AFM) to measure the Young’s modulus of native SEI and correlated the mechanical properties of the SEI with battery performance by introducing maximum elastic strain energy U related with Young’s modulus E and elastic strain limit ɛY of the SEI.

Gao et al. developed an AFM-based nanoindentation test to separately measure the Young’s modulus and elastic strain limit of SEI. In Fig. 13(a), the AFM probe exerts a force on the SEI sample in the elastic stage to the maximum load, and then lifts back to the original point 1. The Young’s modulus of the SEI is obtained by fitting the force–displacement unloading curve between points 2 and 3, and point 2 represents the point where the deformation of the sample begins. In the second step from Fig. 13(b), the AFM probe directly applies a large load to the sample to make it fail, aimed at acquiring the data related to its elastic strain limit. From the relevant force–displacement curve, we find that point a could be experimentally recognized as the first deviation point of the fitting curve of the Hertz contact model. The curve from point b to point c implies that the sample under the probe will appear plastic deformation and even break. Then, the probe continues to apply the load until the sample fails completely at point d and the resulting curve from point c deviates from the fitting curve of the Hertz contact model. It is worth noting that point b is a critical point to determine the elastic strain limit of the SEI since it is the first discontinuity point during the plastic yield stage. Through the two-step measurement by AFM, the Young’s modulus of the native SEI formed in various electrolytes ranges from 0.3 GPa to 1.2 GPa, which is summarized in Fig. 13(c). Gao et al. used the finite element method to simulate the deformation distribution of SEI under the external pressure applied by the probe. It is found that the deformed region extends to the anode beyond the SEI, showing the non-negligible influence of the anode on the SEI. Therefore, errors could exist in measuring the Young’s modulus of SEI if the effect of the anode is not considered.

Fig. 13
Fig. 13
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Furthermore, Gao et al. introduced a mechanical parameter of maximum elastic strain energy U, defined by the following formula:
$U=815(45π)5⋅r3⋅(1−ν2)4⋅E⋅(εY)5$
where ν is the SEI Poisson’s ratio, and r is the radius of the rigid indenter. And the values of E and ɛY are obtained through the above AFM experiments. U reflects the stability of SEI, which is embodied in the Coulombic efficiency of the battery.

In Fig. 14, the average Coulombic efficiency (ACE) of the battery is basically proportional to the U of the SEI. Besides, the SEI with a relatively small U easily fractures during Li plating. This is mainly because the energy generated by Li metal cannot be completely absorbed by the SEI with a small U, which causes excess energy to be consumed only by the rupture of the SEI, resulting in the instability and low CE of the battery. Hence, the ability of SEI to withstand fracture increases with increasing U. The high U value will promise the strength of the SEI, thereby improving the ACE of the battery as displayed in the 1 M KFSI/DME system. In short, the proposal of U is quite helpful for how the mechanical properties of the SEI affect the performance of the battery.

Fig. 14
Fig. 14
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In addition to analyzing the influence of maximum elastic strain energy U on the battery performance, Gao et al. showed that the native SEI possesses poor mechanical properties. In general, the modulus of inorganic components in SEI is much larger than that of organic ones. Gao et al. simulated the effect of inorganic components on the mechanical parameters of the whole SEI through the finite element method. The simulation results showed that both the maximum elastic strain energy U and Young’s modulus E of the entire SEI are significantly small, caused by the large mismatch in the Young’s moduli between organic components and inorganic components, which makes the SEI be prone to yield. In addition, the complexity of SEI components results in unstable mechanical properties, and thus artificial SEI contains a single component or a combination of a few components.

The AFM is a common tool for measuring the Young’s modulus of SEI. Liu et al. [53] used AFM to measure the Young’s modulus of native SEI in conventional batteries and found it to be in the range of ~100 MPa. Zhang et al. [54] also measured the Young’s modulus and thickness of SEI in the specific mixed electrolytes by ATM and analyzed their relationship.

It is worth noting that the native SEI has a double-layer structure containing organic–inorganic components. The coexistence of the bilayer structure further confirms the inhomogeneity of the native SEI. Figure 15(a) illustrates that for the SEI with tens of nanometers, the Young’s modulus is mainly in the range of 100–1000 MPa, and there is a corresponding variation with the change of thickness. Figure 15(b) shows the distribution of Young’s modulus within the SEI, concentrated between 25 and 27 (10 MPa). Thus, the distribution of Young’s modulus with SEI thickness and composition is highly heterogeneous.

Fig. 15
Fig. 15
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The Young’s modulus of SEI could determine whether the SEI can suppress Li dendrite. The native SEI layer has Young’s modulus of 2.5 GPa in 1 M LiTFSI DOL: DME electrolyte and 1.7 GPa in carbonate electrolytes [55,56]. However, the modulus of SEI formed in electrolytes with different concentrations and species can be completely varied. For example, the moduli of SEI correspond to 1.3 GPa, 1.6 GPa, and 1.0 GPa in no-salt P14TFSI, 1 M LiTFSI in P14TFSI, and 1 M LiFSI in P14TFSI, respectively [57,58]. Therefore, the native SEI tends to have low Young’s modulus, and thus, it is not capable of suppressing the growth of Li dendrite. In addition, the porosity and structural variety of SEI decrease the overall Young’s modulus. Shen et al. [36] found that SEI with low modulus can cause nonuniform deposition. It illustrates that the SEI with poor mechanical performance could not buffer the volume changes of Li metal anode, which promotes dendrite growth. Moreover, they also provided an ideal value of Young’s modulus, 3 GPa, to withstand the volume changes of Li metal anode and suppress Li dendrite growth.

## 3 Strategies of Regulating Solid Electrolyte Interphase to Suppress Lithium Dendrite

As mentioned above, native SEI has disadvantages, among which low ionic conductivity, inhomogeneity, and poor mechanical properties are the main causes of promoting Li dendrite growth. In recent years, researchers have made a lot of efforts to improve the SEI stability, such as electrolyte additives and artificial SEI. The purpose is to obtain a stable and robust SEI to suppress the growth of dendrite and improve battery cycle life. In this section, we discuss some specific methods to achieve ideal SEI characters for uniform Li electrodeposition.

For conventional EC electrolytes, cracks appear in SEI, which is difficult to restore to the original SEI morphology. After adding FEC, only a few cracks are observed even with the strain up to 6.2%, further illustrating that SEI in EC with FEC electrolyte has enhanced the mechanical resistance to break (Fig. 16). DMS can also act as a kind of additive [62]. The SEI under DMS additive has large Young’s modulus, compact structure, and high ionic conductivity. Besides, it has an electrochemically stable impedance with a low value of 10 Ω.

Fig. 16
Fig. 16
Close modal

Fig. 17
Fig. 17
Close modal

### 3.2 Electrolyte Composition.

It is possible to modify the SEI by appropriately changing the composition ratio of electrolytes. Fan et al. [67] mixed three common electrolyte components in certain proportions to create a new all-fluorinated electrolyte. It contains 1 M LiPF6 with a mixture of fluoroethylene carbonate/2,2,2-trifluoroethyl methyl carbonate/1,1,2,2-tetrafluoroethyl 2′,2′,2′-trifluoroethyl ether (FEC: FEMC: HFE, 2:6:2 by wt.), which has 22 M F and easily form SEI with rich LiF inorganize component (Fig. 18).

Fig. 18
Fig. 18
Close modal

By varying the electrolyte composition ratio, Fan et al. acquired the SEI having an extremely high LiF component proportion (about 45%), and the SEI could suppress the growth of Li dendrite. Wu et al. [68] adjusted the ratio of the two electrolyte components (bis (trifluoromethanesulfonyl)-imide/TEGDME electrolyte and LiTFSI/lithium difluoro(oxalate)borate (LiODFB)) to 6:4 in Li–S battery. Subsequent SEI has excellent strength to protect the cell from undergoing Li dendrite. In addition to changing the electrolyte composition ratio, increasing the salt concentration in the electrolyte is also a common and effective method. The native SEI formed in 1 M LiTFSI in 1:1 DOL: DME has Young’s modulus of 0.3 GPa, while the modulus of the SEI formed under 7 M LiTFSI in 1:1 DOL: DME electrolyte increases to 2.8 GPa [69]. Wang et al. [70] found that the SEI formed under 1 M LiFSI in DME electrolyte has a modulus of 3.5 GPa. They increased the salt concentration of the electrolyte to 2 M and obtained a SEI with significantly increased modulus of 10.7 GPa. However, as the electrolyte salt concentration increased to 3 M, the modulus of the SEI decreased to 4.2 GPa. The cell showed better performance in the 2 M concentration electrolyte than those in 1 M and 3 M.

Cui et al. [71] employed Li|Li symmetric cells with 1 M and 8 M LiTFSI-PC electrolytes to explore the effect of electrolyte concentration on cell performance. They found that the SEI layer fractured in 1 M electrolyte after only 10 cycles, resulting in severe structural damage during subsequent cycles. In contrast, the battery can be maintained with a much longer time in 8 M electrolyte, due to the excellent electrochemical stability of the SEI. Besides, Cui et al. also used Cu|Li system with 1 M and 8 M LiTFSI-PC electrolytes to observe the variation in SEI morphology. In Fig. 19, the SEI layer formed in 1 M LiTFSI-PC electrolyte shows an agglomerated structure with severe cracks and loose morphology. It may be due to the incompatibility between Li metal and electrolyte. For the 8 M electrolyte, the SEI exhibits tough, compact, dense, and uninterrupted morphology, which confirms that increasing salt concentration helps suppress dendrite and prolong battery life.

Fig. 19
Fig. 19
Close modal

### 3.3 Artificial Solid Electrolyte Interphase.

Adding electrolyte additives and changing electrolyte composition are typical methods for modifying electrolytes to obtain stable SEI. Electrolyte additives can maintain short-term stable cycles because trace additives are easily consumed by reduction reactions during Li plating [7274]. For high concentration electrolytes, the salt concentration is sufficient to maintain the battery for long-term cycling without failure. Furthermore, the high concentration of solvent molecules can effectively prevent further side reactions between Li metal and electrolyte. However, using high concentration electrolytes increases the cost. And the electrolyte becomes more viscous as the concentration increases, which is detrimental to the transport of Li ions [7577]. Different from modifying electrolytes, artificial SEI could reduce the consumption of Li metal and electrolytes and eliminate the issues of Li dendrite.

Guo et al. [78] fabricated uniform Li3PO4 artificial SEI film by in situ reaction of polyphosphoric acid (PPA) and Li metal. In Fig. 20, the inhomogeneous and fragile pristine SEI can cause irregular Li deposition and repetitive fracture regeneration of the film, leading to the formation of dendrite. The performance of the cell containing Li with pristine SEI is greatly degraded as the number of cycles increases (Fig. 20(d)). Replacing the pristine SEI with the uniform Li3PO4 artificial SEI leads to dendrite-free Li morphology (Fig. 20(b)). Moreover, Li3PO4 artificial SEI has Young’s modulus of 10–11 GPa, which can be sufficient to inhibit the growth of dendrite and prevent side reactions between Li metal and electrolytes. The cell with this LiPPA-based SEI also exhibits superior cycling performance (Fig. 20(d)). Zhou et al. [79] designed a novel organophosphorus hybrid LixPO4 artificial SEI layer for PO4 group-containing species to ensure dendrite-free Li morphology and high-performance cycling.

Fig. 20
Fig. 20
Close modal

In Figs. 21(a)21(d), organophosphorus hybrid SEI (OPHS) film provides a compact and dense structure. The hybrid artificial SEI can offer uniform deposition channels for Li ions, ensuring dendrite-free morphology on the anode surface during electrodeposition. The Cu||Li cells with OPHS-Li anode can maintain Coulombic efficiency up to 99% at the current density of 0.5 mA cm−2 even after 500 cycles, while the Coulombic efficiency of Cu||Li cells without OPHS decreases to 80% after 150 cycles (Fig. 21(e)). As the current density increases to 4.0 mA cm−2, the cells containing OPHS can still maintain Coulombic efficiency of 97% after 100 cycles. In contrast, the ones with bare Li exhibit unstable cycling performance, and the Coulombic efficiency decreases to 70% after 40 cycles (Fig. 21(f)).

Fig. 21
Fig. 21
Close modal

Zhang et al. [80] designed a bilayer hybrid artificial SEI film containing organic species PEO and Li alkylcarbonate and inorganic species Li salts such as LiF and Li3N. Unlike the single-component SEI, this artificial SEI possesses both the flexibility of organics and the robustness of inorganics. Moreover, it can combine the advantages of the two components to form a synergistic effect, which facilitates uniform ion transport and suppress Li dendrite growth (Fig. 22(a)). Xu et al. [81] also designed a double-layer hybrid artificial SEI film, which is composed of the soft organics PVDF-HFP and robust inorganics LiF. In Fig. 22(b), the Li metal morphologies become extremely rough on the bare Li and single PVDF-HFP coated Li anode. In contrast, the constructed hybrid artificial SEI has high ionic conductivity, superior mechanical strength, and suitable uniformity. Thus, Li deposition morphology is uniform and smooth on the protective layer coated Li anode. Kwak et al. [82] designed a Nafion-Al2O3 hybrid artificial SEI for Li-O2 cells. Notably, Nafion has high ionic conductivity and excellent electrochemical stability, which satisfy the characteristics of ideal SEI. However, the mechanical properties of Nafion are too weak to suppress the growth of dendrite. Kwak et al. proposed to add inorganic components with superior mechanical properties, such as AI2O3 added to the original Nafion to exert a synergistic effect. In Fig. 22(c), the cells with Nafion-Al2O3 hybrid film exhibit uniform morphology, while dendrites and dead Li are generated inside the ones with Nafion alone. Therefore, the hybrid film can effectively suppress dendritic Li growth and improve the cycling performance of batteries due to its superior synergistic properties.

Fig. 22
Fig. 22
Close modal

Kim et al. [83] designed a PEDOT-co-PEG polymer artificial SEI to improve battery cycling performance. The copolymer has not only high ionic conductivity but also strong adhesion to Li metal anode, which effectively suppresses dendritic Li growth. Ma et al. [84] prepared the PEDOT-co-PEG copolymer film for Li–S batteries to test cycling performance (Fig. 23(a)). The results show that the cells with copolymer film exhibit relatively high and stable capacity retention, and the Coulombic efficiency remains nearly unchanged after 300 cycles, while the ones with native SEI show an obvious decreasing trend (Fig. 23(b)). It is attributed to the PEDOT-co-PEG copolymer artificial SEI that can promote uniform Li deposition and suppress the growth of Li dendrite due to its high ionic conductivity and mechanically strong adhesion. The cycling performance of the batteries can thus be significantly improved. Dong et al. [85] designed an 18-Crown-6 and PVDF composite polymeric artificial SEI coated on the anode through spin-coating method, with a thickness of 5.1 μm (Fig. 23(c)). The composite artificial SEI also has synergistic advantages, including high electrochemical stability of PVDF and high ionic conductivity of 18-Crown-6, which help achieve uniform deposition and suppress dendritic Li growth.

Fig. 23
Fig. 23
Close modal

Zheng et al. [86] designed a monolayer nanostructured artificial SEI, which contains multiple amorphous hollow carbon nanospheres, as shown in Fig. 24. As novel nanomaterials, hollow carbon artificial SEI has an ultrahigh ionic conductivity of 7.5 S m−1 to facilitate the transport of Li ions and extremely high Young’s modulus of 200 GPa to suppress dendrite. Besides, the flexibility and compactness of carbon material accommodate the volume change of Li metal and eliminate the side reactions between the anode and electrolytes.

Fig. 24
Fig. 24
Close modal

Gao et al. [87] designed a skin-grafting artificial SEI using the poly ((N-2,2-dimethyl-1,3-dioxolane-4-methyl)-5-norbornene-exo-2,3-dicarboximide). The grafted skin artificial SEI has a high volume fraction of polymer chains cyclic ether groups that can adhere tightly to the surface of Li metal and prevent detrimental reactions (Fig. 25). Furthermore, it can also regulate the Li+ plating/stripping behavior and eliminate dendrite formation, thereby improving battery performance with higher CE and better cycling stability.

Fig. 25
Fig. 25
Close modal

To sum up, most artificial SEIs have the characteristics of the ideal SEIs, such as high ionic conductivity, homogeneity, superior mechanical properties, and high electrochemical stability, all of which are crucial factors for suppressing dendrite growth and improving battery performance. In addition, copolymerization or composition with other organic/inorganic components is a common method to construct hybrid artificial SEIs to achieve improved mechanical properties and flexibility of SEIs and eliminate dendrites through the synergistic effect between components. The hybridization via polymerization exhibits unique advantages in improving battery properties. Moreover, some artificial SEI films provide preferential diffusion channels for Li ions and prevent electrons from transporting through the films, enabling fast Li plating/stripping to facilitate deposition. The interactions of artificial SEI films with Li metal surfaces, such as enhanced adhesion, not only alleviate side reactions between electrolyte and Li metal but also play an important role in suppressing dendrite growth. Although great progress has been made in fabricating artificial SEI films on Li metal anode, the majority of these techniques have not been used in practical applications. At present, a comprehensive understanding of ideal artificial SEIs, including their compositions and properties, is comparatively lacking. Therefore, the design concepts of artificial SEI are highly limited, and achieving long-term stable cycles of LMBs is desired [88]. Currently, artificial SEIs could not completely exert the synergistic effect of various components such as organics and inorganics, which is attributed to the insufficient fundamental research [89]. In high-energy-density LMBs, Li ions deposit during cycles, leading to the dendritic growth. However, artificial SEIs are not smart and adaptive to the Li metal anode. Therefore, the cracks caused by the volume change could also occur on the artificial SEI and accelerate the inhomogeneous Li deposition. The cells with artificial SEIs could exhibit worse cycling stability and increased polarization, resulting in the rapid deterioration of electrochemical performance [90,91]. In addition to the good properties of artificial SEI, its matching with ultrahigh-energy-capacity cells is also essential. Further studies on designing smart, adaptive, and stable artificial SEIs are necessary to protect Li anode in the next-generation commercial batteries.

## 4 Conclusion

In this review, the underlying mechanisms of SEI–dendrite interactions are discussed, and the corresponding strategies toward suppressing Li dendrite growth and improving battery performance are summarized. Li electrodeposition morphology strongly depends on the SEI properties on the anode surface. However, the native SEI generated from spontaneous reactions between electrolytes and Li metal is electrochemically and mechanically unstable. Low ionic conductivity induces a large ion concentration gradient within the SEI film, and Li ions are preferentially deposited at the tip of Li nucleus, causing the deposited Li metal to grow vertically rather than horizontally. Homogeneous ion transport is also hindered by inhomogeneities in the structure and composition of the native SEI. The SEI structure largely determines the electric potential gradient within the film and Li deposition behavior, and the SEI composition determines ionic conductivity and mechanical robustness [92]. Native SEI is composed of various organic and inorganic components, and its thickness is spatially varied. Those results in different deposition rates on the Li anode, promoting nonuniform deposition and dendrite formation. In addition, the mechanical properties of native SEI are weak. The effective Young’s modulus of SEI is several orders of magnitude lower than single inorganic component due to the large mismatch between organics and inorganics inside the SEI. Thus, the SEI could not withstand the infinite volume change of the anode during Li plating. If Li dendrite forms, native SEI subsequently undergoes repeated fracture and regeneration processes, resulting in the consumption of active materials and degraded battery performance. Fortunately, increasing efforts have been made to modify the SEI, such as modification of electrolytes and artificial SEI. The modified SEI can effectively suppress and even eliminate the growth of Li dendrite due to its ideal properties, i.e., high ionic conductivity, superior mechanics, and electrochemical stability. For high-energy-density batteries, Li metal and silicon are the next-generation anodes, which have the potential to replace the commercial graphite anode. The native SEI on Li metal anode possesses some common species with those on the graphite and Si anode, illustrating that SEIs on these anodes have similar structures and properties such as inhomogeneity and Young’s modulus. Similar to the Li metal anode, the large volume change of Si anode also causes SEI failure and poor cycle performance [93]. Therefore, Li metal–SEI interactions mentioned above could be applicable to the Si system. However, the understanding and strategies for stable Li metal—SEI interactions are not completely applicable to other metal anodes, such as Na, K, Mg, and Ca. Among these metals, Na and K exhibit lower non-negativity, resulting in different properties of SEI components. And the hardness of Na/K is lower than that of Li, showing that it may be easier to modify SEI to suppress dendritic growth. In contrast, multivalent metals, such as Mg and Ca, have high moduli, and the growth of dendrites becomes more dangerous. Furthermore, the high charge density of multivalent ions leads to lower ionic conductivity of SEI, which increases the difficulties of regulating SEI [94].

## Acknowledgment

The information, data, or work presented herein were funded by the Natural Science Foundation of China (Grant No. 12002192) and the Natural Science Foundation of Shandong Province, China (Grant No. ZR2020QA043).

## Conflict of Interest

There are no conflicts of interest.

## References

1.
Grande
,
L.
,
Paillard
,
E.
,
Hassoun
,
J.
,
Park
,
J. B.
,
Lee
,
Y. J.
,
Sun
,
Y. K.
,
Passerini
,
S.
, and
Scrosati
,
B.
,
2015
, “
The Lithium/Air Battery: Still an Emerging System or a Practical Reality?
,”
,
27
(
5
), pp.
784
800
.
2.
Lu
,
Y.
,
Tikekar
,
M.
,
Mohanty
,
R.
,
Hendrickson
,
K.
,
Ma
,
L.
, and
Archer
,
L. A.
,
2015
, “
Stable Cycling of Lithium Metal Batteries Using High Transference Number Electrolytes
,”
,
5
(
9
), p.
1402073
.
3.
Xu
,
W.
,
Wang
,
J.
,
Ding
,
F.
,
Chen
,
X.
,
Nasybulin
,
E.
,
Zhang
,
Y.
, and
Zhang
,
J.-G.
,
2014
, “
Lithium Metal Anodes for Rechargeable Batteries
,”
Energy Environ. Sci.
,
7
(
2
), pp.
513
537
.
4.
Cheng
,
X.-B.
, and
Zhang
,
Q.
,
2015
, “
Dendrite-Free Lithium Metal Anodes: Stable Solid Electrolyte Interphases for High-Efficiency Batteries
,”
J. Mater. Chem. A
,
3
(
14
), pp.
7207
7209
.
5.
Sun
,
Y.
,
Liu
,
N.
, and
Cui
,
Y.
,
2016
, “
Promises and Challenges of Nanomaterials for Lithium-Based Rechargeable Batteries
,”
Nat. Energy
,
1
(
7
), p.
16071
.
6.
Tikekar
,
M. D.
,
Choudhury
,
S.
,
Tu
,
Z.
, and
Archer
,
L. A.
,
2016
, “
Design Principles for Electrolytes and Interfaces for Stable Lithium-Metal Batteries
,”
Nat. Energy
,
1
(
9
), p.
16114
.
7.
Xie
,
J.
,
Wang
,
J.
,
Lee
,
H. R.
,
Yan
,
K.
,
Li
,
Y.
,
Shi
,
F.
,
Huang
,
W.
,
Pei
,
A.
,
Chen
,
G.
, and
Subbaraman
,
R.
,
2018
, “
Engineering Stable Interfaces for Three-Dimensional Lithium Metal Anodes
,”
,
4
(
7
), p.
eaat5168
.
8.
Xu
,
X.
,
Wang
,
S.
,
Wang
,
H.
,
Hu
,
C.
,
Jin
,
Y.
,
Liu
,
J.
, and
Yan
,
H.
,
2018
, “
Recent Progresses in the Suppression Method Based on the Growth Mechanism of Lithium Dendrite
,”
J. Energy Chem.
,
27
(
2
), pp.
513
527
.
9.
Cheng
,
X. B.
,
Zhang
,
R.
,
Zhao
,
C. Z.
, and
Zhang
,
Q.
,
2017
, “
Toward Safe Lithium Metal Anode in Rechargeable Batteries: A Review
,”
Chem. Rev.
,
117
(
15
), pp.
10403
10473
.
10.
Fang
,
C.
,
Li
,
J.
,
Zhang
,
M.
,
Zhang
,
Y.
,
Yang
,
F.
,
Lee
,
J. Z.
,
Lee
,
M. H.
, et al
,
2019
, “
Quantifying Inactive Lithium in Lithium Metal Batteries
,”
Nature
,
572
(
7770
), pp.
511
515
.
11.
Dornbusch
,
D. A.
,
Hilton
,
R.
,
Lohman
,
S. D.
, and
Suppes
,
G. J.
,
2014
, “
Experimental Validation of the Elimination of Dendrite Short-Circuit Failure in Secondary Lithium-Metal Convection Cell Batteries
,”
J. Electrochem. Soc.
,
162
(
3
), pp.
A262
A268
.
12.
Li
,
W.
,
Zheng
,
H.
,
Chu
,
G.
,
Luo
,
F.
,
Zheng
,
J.
,
Xiao
,
D.
,
Li
,
X.
, et al
,
2014
, “
Effect of Electrochemical Dissolution and Deposition Order on Lithium Dendrite Formation: A Top View Investigation
,”
,
176
, pp.
109
124
.
13.
Cohen
,
Y. S.
,
Cohen
,
Y.
, and
Aurbach
,
D.
,
2000
, “
Micromorphological Studies of Lithium Electrodes in Alkyl Carbonate Solutions Using In Situ Atomic Force Microscopy
,”
J. Phys. Chem. B
,
104
(
51
), pp.
12282
12291
.
14.
Cheng
,
X. B.
,
Zhang
,
R.
,
Zhao
,
C. Z.
,
Wei
,
F.
,
Zhang
,
J. G.
, and
Zhang
,
Q.
,
2016
, “
A Review of Solid Electrolyte Interphases on Lithium Metal Anode
,”
,
3
(
3
), p.
1500213
.
15.
Peled
,
E.
,
Golodnitsky
,
D.
, and
Ardel
,
G.
,
1997
, “
Advanced Model for Solid Electrolyte Interphase Electrodes in Liquid and Polymer Electrolytes
,”
J. Electrochem. Soc.
,
144
(
8
), pp.
L208
L210
.
16.
Christensen
,
J.
, and
Newman
,
J.
,
2004
, “
A Mathematical Model for the Lithium-Ion Negative Electrode Solid Electrolyte Interphase
,”
J. Electrochem. Soc.
,
151
(
11
), p.
A1977
.
17.
Borodin
,
O.
,
Smith
,
G. D.
, and
Fan
,
P.
,
2006
, “
Molecular Dynamics Simulations of Lithium Alkyl Carbonates
,”
J. Phys. Chem. B
,
110
(
45
), pp.
22773
22779
.
18.
Ramasubramanian
,
A.
,
Yurkiv
,
V.
,
Foroozan
,
T.
,
Ragone
,
M.
,
Shahbazian-Yassar
,
R.
, and
Mashayek
,
F.
,
2019
, “
Lithium Diffusion Mechanism Through Solid–Electrolyte Interphase in Rechargeable Lithium Batteries
,”
J. Phys. Chem. C
,
123
(
16
), pp.
10237
10245
.
19.
Shi
,
S.
,
Lu
,
P.
,
Liu
,
Z.
,
Qi
,
Y.
,
Hector
,
L. G.
, Jr.
,
Li
,
H.
, and
Harris
,
S. J.
,
2012
, “
Direct Calculation of Li-Ion Transport in the Solid Electrolyte Interphase
,”
J. Am. Chem. Soc.
,
134
(
37
), pp.
15476
15487
.
20.
Single
,
F.
,
Horstmann
,
B.
, and
Latz
,
A.
,
2017
, “
Revealing SEI Morphology: In-Depth Analysis of a Modeling Approach
,”
J. Electrochem. Soc.
,
164
(
11
), pp.
E3132
E3145
.
21.
Yildirim
,
H.
,
Kinaci
,
A.
,
Chan
,
M. K.
, and
Greeley
,
J. P.
,
2015
, “
First-Principles Analysis of Defect Thermodynamics and Ion Transport in Inorganic SEI Compounds: LiF and NaF
,”
ACS Appl. Mater. Interfaces
,
7
(
34
), pp.
18985
18996
.
22.
Mulliner
,
A. D.
,
Aeberhard
,
P. C.
,
Battle
,
P. D.
,
David
,
W. I.
, and
Refson
,
K.
,
2015
, “
Diffusion in Li(2)O Studied by Non-equilibrium Molecular Dynamics for 873 < T/K < 1603
,”
Phys. Chem. Chem. Phys.
,
17
(
33
), pp.
21470
21475
.
23.
Wu
,
H.
,
Jia
,
H.
,
Wang
,
C.
,
Zhang
,
J. G.
, and
Xu
,
W.
,
2021
, “
Recent Progress in Understanding Solid Electrolyte Interphase on Lithium Metal Anodes
,”
,
11
(
5
), p.
2003092
.
24.
Chen
,
Y. C.
,
Ouyang
,
C. Y.
,
Song
,
L. J.
, and
Sun
,
Z. L.
,
2011
, “
Electrical and Lithium Ion Dynamics in Three Main Components of Solid Electrolyte Interphase From Density Functional Theory Study
,”
J. Phys. Chem. C
,
115
(
14
), pp.
7044
7049
.
25.
Pan
,
J.
,
Cheng
,
Y.-T.
, and
Qi
,
Y.
,
2015
, “
General Method to Predict Voltage-Dependent Ionic Conduction in a Solid Electrolyte Coating on Electrodes
,”
Phys. Rev. B
,
91
(
13
), p.
134116
.
26.
Borodin
,
O.
,
Zhuang
,
G. V.
,
Ross
,
P. N.
, and
Xu
,
K.
,
2013
, “
Molecular Dynamics Simulations and Experimental Study of Lithium Ion Transport in Dilithium Ethylene Dicarbonate
,”
J. Phys. Chem. C
,
117
(
15
), pp.
7433
7444
.
27.
Schafzahl
,
L.
,
Ehmann
,
H.
,
Kriechbaum
,
M.
,
Sattelkow
,
J.
,
Ganner
,
T.
,
Plank
,
H.
,
Wilkening
,
M.
, and
Freunberger
,
S. A.
,
2018
, “
Long-Chain Li and Na Alkyl Carbonates as Solid Electrolyte Interphase Components: Structure, Ion Transport, and Mechanical Properties
,”
Chem. Mater.
,
30
(
10
), pp.
3338
3345
.
28.
Wang
,
L.
,
Menakath
,
A.
,
Han
,
F.
,
Wang
,
Y.
,
Zavalij
,
P. Y.
,
,
K. J.
,
Borodin
,
O.
, et al
,
2019
, “
Identifying the Components of the Solid-Electrolyte Interphase in Li-Ion Batteries
,”
Nat. Chem.
,
11
(
9
), pp.
789
796
.
29.
Lorger
,
S.
,
Usiskin
,
R.
, and
Maier
,
J.
,
2019
, “
Transport and Charge Carrier Chemistry in Lithium Oxide
,”
J. Electrochem. Soc.
,
166
(
10
), pp.
A2215
A2220
.
30.
Guo
,
R.
, and
Gallant
,
B. M.
,
2020
, “
Li2O Solid Electrolyte Interphase: Probing Transport Properties at the Chemical Potential of Lithium
,”
Chem. Mater.
,
32
(
13
), pp.
5525
5533
.
31.
Shi
,
S.
,
Qi
,
Y.
,
Li
,
H.
, and
Hector
,
L. G.
,
2013
, “
Defect Thermodynamics and Diffusion Mechanisms in Li2CO3 and Implications for the Solid Electrolyte Interphase in Li-Ion Batteries
,”
J. Phys. Chem. C
,
117
(
17
), pp.
8579
8593
.
32.
Lowe
,
J. S.
, and
Siegel
,
D. J.
,
2020
, “
Modeling the Interface Between Lithium Metal and Its Native Oxide
,”
ACS Appl. Mater. Interfaces
,
12
(
41
), pp.
46015
46026
.
33.
Vu
,
T. T.
,
Eom
,
G. H.
,
Lee
,
J.
,
Park
,
M.-S.
, and
Moon
,
J.
,
2021
, “
Electrolyte Interface Design for Regulating Li Dendrite Growth in Rechargeable Li-Metal Batteries: A Theoretical Study
,”
J. Power Sources
,
496
, p.
229791
.
34.
Yoon
,
G.
,
Moon
,
S.
,
Ceder
,
G.
, and
Kang
,
K.
,
2018
, “
Deposition and Stripping Behavior of Lithium Metal in Electrochemical System: Continuum Mechanics Study
,”
Chem. Mater.
,
30
(
19
), pp.
6769
6776
.
35.
Veith
,
G. M.
,
Doucet
,
M.
,
Sacci
,
R. L.
,
Vacaliuc
,
B.
,
Baldwin
,
J. K.
, and
Browning
,
J. F.
,
2017
, “
Determination of the Solid Electrolyte Interphase Structure Grown on a Silicon Electrode Using a Fluoroethylene Carbonate Additive
,”
Sci. Rep.
,
7
(
1
), p.
6326
.
36.
Shen
,
X.
,
Zhang
,
R.
,
Chen
,
X.
,
Cheng
,
X. B.
,
Li
,
X.
, and
Zhang
,
Q.
,
2020
, “
The Failure of Solid Electrolyte Interphase on Li Metal Anode: Structural Uniformity or Mechanical Strength?
,”
,
10
(
10
), p.
1903645
.
37.
He
,
X.
,
Bresser
,
D.
,
Passerini
,
S.
,
Baakes
,
F.
,
Krewer
,
U.
,
Lopez
,
J.
,
Mallia
,
C. T.
, et al
,
2021
, “
The Passivity of Lithium Electrodes in Liquid Electrolytes for Secondary Batteries
,”
Nat. Rev. Mater.
,
6
(
11
), pp.
1036
1052
.
38.
Liu
,
G.
, and
Lu
,
W.
,
2017
, “
A Model of Concurrent Lithium Dendrite Growth, SEI Growth, SEI Penetration and Regrowth
,”
J. Electrochem. Soc.
,
164
(
9
), pp.
A1826
A1833
.
39.
Dey
,
A.
, and
Sullivan
,
B.
,
1970
, “
The Electrochemical Decomposition of Propylene Carbonate on Graphite
,”
J. Electrochem. Soc.
,
117
(
2
), p.
222
.
40.
Peled
,
E.
,
1979
, “
The Electrochemical Behavior of Alkali and Alkaline Earth Metals in Nonaqueous Battery Systems—The Solid Electrolyte Interphase Model
,”
J. Electrochem. Soc.
,
126
(
12
), pp.
2047
2051
.
41.
Aurbach
,
D.
,
2000
, “
Review of Selected Electrode–Solution Interactions Which Determine the Performance of Li and Li Ion Batteries
,”
J. Power Sources
,
89
(
2
), pp.
206
218
.
42.
Sacci
,
R. L.
,
Black
,
J. M.
,
Balke
,
N.
,
Dudney
,
N. J.
,
More
,
K. L.
, and
Unocic
,
R. R.
,
2015
, “
Nanoscale Imaging of Fundamental Li Battery Chemistry: Solid-Electrolyte Interphase Formation and Preferential Growth of Lithium Metal Nanoclusters
,”
Nano Lett.
,
15
(
3
), pp.
2011
2018
.
43.
Lin
,
Y.-X.
,
Liu
,
Z.
,
Leung
,
K.
,
Chen
,
L.-Q.
,
Lu
,
P.
, and
Qi
,
Y.
,
2016
, “
Connecting the Irreversible Capacity Loss in Li-Ion Batteries With the Electronic Insulating Properties of Solid Electrolyte Interphase (SEI) Components
,”
J. Power Sources
,
309
, pp.
221
230
.
44.
Ozhabes
,
Y.
,
Gunceler
,
D.
, and
Arias
,
T.
,
2015
, “
Stability and surface diffusion at lithium-electrolyte interphases with connections to dendrite suppression
,” arXiv preprint arXiv:1504.05799.
45.
Hao
,
F.
,
Verma
,
A.
, and
Mukherjee
,
P. P.
,
2018
, “
Mechanistic Insight Into Dendrite–SEI Interactions for Lithium Metal Electrodes
,”
J. Mater. Chem. A
,
6
(
40
), pp.
19664
19671
.
46.
Wang
,
X.
,
Zeng
,
W.
,
Hong
,
L.
,
Xu
,
W.
,
Yang
,
H.
,
Wang
,
F.
,
Duan
,
H.
,
Tang
,
M.
, and
Jiang
,
H.
,
2018
, “
Stress-Driven Lithium Dendrite Growth Mechanism and Dendrite Mitigation by Electroplating on Soft Substrates
,”
Nat. Energy
,
3
(
3
), pp.
227
235
.
47.
Yu
,
Z.
,
Cui
,
Y.
, and
Bao
,
Z.
,
2020
, “
Design Principles of Artificial Solid Electrolyte Interphases for Lithium-Metal Anodes
,”
Cell Rep. Phys. Sci.
,
1
(
7
), p.
100119
.
48.
Liu
,
Z.
,
Qi
,
Y.
,
Lin
,
Y. X.
,
Chen
,
L.
,
Lu
,
P.
, and
Chen
,
L. Q.
,
2016
, “
Interfacial Study on Solid Electrolyte Interphase at Li Metal Anode: Implication for Li Dendrite Growth
,”
J. Electrochem. Soc.
,
163
(
3
), pp.
A592
A598
.
49.
Deng
,
Z.
,
Wang
,
Z.
,
Chu
,
I.-H.
,
Luo
,
J.
, and
Ong
,
S. P.
,
2015
, “
Elastic Properties of Alkali Superionic Conductor Electrolytes From First Principles Calculations
,”
J. Electrochem. Soc.
,
163
(
2
), pp.
A67
A74
.
50.
Monroe
,
C.
, and
Newman
,
J.
,
2005
, “
The Impact of Elastic Deformation on Deposition Kinetics at Lithium/Polymer Interfaces
,”
J. Electrochem. Soc.
,
152
(
2
), p.
A396
.
51.
Stone
,
G. M.
,
Mullin
,
S. A.
,
Teran
,
A. A.
,
Hallinan
,
D. T.
,
Minor
,
A. M.
,
Hexemer
,
A.
, and
Balsara
,
N. P.
,
2012
, “
Resolution of the Modulus Versus Adhesion Dilemma in Solid Polymer Electrolytes for Rechargeable Lithium Metal Batteries
,”
J. Electrochem. Soc.
,
159
(
3
), pp.
A222
A227
.
52.
Gao
,
Y.
,
Du
,
X.
,
Hou
,
Z.
,
Shen
,
X.
,
Mai
,
Y.-W.
,
Tarascon
,
J.-M.
, and
Zhang
,
B.
,
2021
, “
Unraveling the Mechanical Origin of Stable Solid Electrolyte Interphase
,”
Joule
,
5
(
7
), pp.
1860
1872
.
53.
Liu
,
X. R.
,
Deng
,
X.
,
Liu
,
R. R.
,
Yan
,
H. J.
,
Guo
,
Y. G.
,
Wang
,
D.
, and
Wan
,
L. J.
,
2014
, “
Single Nanowire Electrode Electrochemistry of Silicon Anode by In Situ Atomic Force Microscopy: Solid Electrolyte Interphase Growth and Mechanical Properties
,”
ACS Appl. Mater. Interfaces
,
6
(
22
), pp.
20317
20323
.
54.
Zhang
,
J.
,
Wang
,
R.
,
Yang
,
X.
,
Lu
,
W.
,
Wu
,
X.
,
Wang
,
X.
,
Li
,
H.
, and
Chen
,
L.
,
2012
, “
Direct Observation of Inhomogeneous Solid Electrolyte Interphase on MnO Anode With Atomic Force Microscopy and Spectroscopy
,”
Nano Lett.
,
12
(
4
), pp.
2153
2157
.
55.
Zheng
,
J.
,
Zheng
,
H.
,
Wang
,
R.
,
Ben
,
L.
,
Lu
,
W.
,
Chen
,
L.
,
Chen
,
L.
, and
Li
,
H.
,
2014
, “
3D Visualization of Inhomogeneous Multi-layered Structure and Young's Modulus of the Solid Electrolyte Interphase (SEI) on Silicon Anodes for Lithium Ion Batteries
,”
Phys. Chem. Chem. Phys.
,
16
(
26
), pp.
13229
13238
.
56.
Yuan
,
S.
,
Weng
,
S.
,
Wang
,
F.
,
Dong
,
X.
,
Wang
,
Y.
,
Wang
,
Z.
,
Shen
,
C.
,
Bao
,
J. L.
,
Wang
,
X.
, and
Xia
,
Y.
,
2021
, “
Revisiting the Designing Criteria of Advanced Solid Electrolyte Interphase on Lithium Metal Anode Under Practical Condition
,”
Nano Energy
,
83
, p.
105847
.
57.
Yoon
,
I.
,
Jurng
,
S.
,
Abraham
,
D. P.
,
Lucht
,
B. L.
, and
Guduru
,
P. R.
,
2018
, “
In Situ Measurement of the Plane-Strain Modulus of the Solid Electrolyte Interphase on Lithium-Metal Anodes in Ionic Liquid Electrolytes
,”
Nano Lett.
,
18
(
9
), pp.
5752
5759
.
58.
Zhang
,
H.
,
Shen
,
C.
,
Huang
,
Y.
, and
Liu
,
Z.
,
2021
, “
Spontaneously Formation of SEI Layers on Lithium Metal From LiFSI/DME and LiTFSI/DME Electrolytes
,”
Appl. Surf. Sci.
,
537
, p.
147983
.
59.
Zhang
,
X.-Q.
,
Cheng
,
X.-B.
,
Chen
,
X.
,
Yan
,
C.
, and
Zhang
,
Q.
,
2017
, “
Fluoroethylene Carbonate Additives to Render Uniform Li Deposits in Lithium Metal Batteries
,”
,
27
(
10
), p.
1605989
.
60.
Chen
,
L.
,
Chen
,
K. S.
,
Chen
,
X.
,
Ramirez
,
G.
,
Huang
,
Z.
,
Geise
,
N. R.
,
Steinruck
,
H. G.
, et al
,
2018
, “
Novel ALD Chemistry Enabled Low-Temperature Synthesis of Lithium Fluoride Coatings for Durable Lithium Anodes
,”
ACS Appl. Mater. Interfaces
,
10
(
32
), pp.
26972
26981
.
61.
Yoon
,
I.
,
Jurng
,
S.
,
Abraham
,
D. P.
,
Lucht
,
B. L.
, and
Guduru
,
P. R.
,
2020
, “
Measurement of Mechanical and Fracture Properties of Solid Electrolyte Interphase on Lithium Metal Anodes in Lithium Ion Batteries
,”
Energy Storage Mater.
,
25
, pp.
296
304
.
62.
Wan
,
G.
,
Guo
,
F.
,
Li
,
H.
,
Cao
,
Y.
,
Ai
,
X.
,
Qian
,
J.
,
Li
,
Y.
, and
Yang
,
H.
,
2018
, “
Suppression of Dendritic Lithium Growth by In Situ Formation of a Chemically Stable and Mechanically Strong Solid Electrolyte Interphase
,”
ACS Appl. Mater. Interfaces
,
10
(
1
), pp.
593
601
.
63.
Gofer
,
Y.
,
Ben-Zion
,
M.
, and
Aurbach
,
D.
,
1992
, “
Solutions of LiAsF6 in 1, 3-Dioxolane for Secondary Lithium Batteries
,”
J. Power Sources
,
39
(
2
), pp.
163
178
.
64.
Eshetu
,
G. G.
,
Judez
,
X.
,
Li
,
C.
,
Bondarchuk
,
O.
,
Rodriguez-Martinez
,
L. M.
,
Zhang
,
H.
, and
Armand
,
M.
,
2017
, “
Lithium Azide As an Electrolyte Additive for All-Solid-State Lithium-Sulfur Batteries
,”
Angew. Chem. Int. Ed.
,
56
(
48
), pp.
15368
15372
.
65.
Qian
,
J.
,
Xu
,
W.
,
Bhattacharya
,
P.
,
Engelhard
,
M.
,
Henderson
,
W. A.
,
Zhang
,
Y.
, and
Zhang
,
J.-G.
,
2015
, “
Dendrite-Free Li Deposition Using Trace-Amounts of Water As an Electrolyte Additive
,”
Nano Energy
,
15
, pp.
135
144
.
66.
Wang
,
E.
,
Dey
,
S.
,
Liu
,
T.
,
Menkin
,
S.
, and
Grey
,
C. P.
,
2020
, “
Effects of Atmospheric Gases on Li Metal Cyclability and Solid-Electrolyte Interphase Formation
,”
ACS Energy Lett.
,
5
(
4
), pp.
1088
1094
.
67.
Fan
,
X.
,
Chen
,
L.
,
Borodin
,
O.
,
Ji
,
X.
,
Chen
,
J.
,
Hou
,
S.
,
Deng
,
T.
, et al
,
2018
, “
Non-flammable Electrolyte Enables Li-Metal Batteries With Aggressive Cathode Chemistries
,”
Nat. Nanotechnol.
,
13
(
8
), pp.
715
722
.
68.
Wu
,
F.
,
Zhu
,
Q.
,
Chen
,
R.
,
Chen
,
N.
,
Chen
,
Y.
,
Ye
,
Y.
,
Qian
,
J.
, and
Li
,
L.
,
2015
, “
Ionic Liquid-Based Electrolyte With Binary Lithium Salts for High Performance Lithium–Sulfur Batteries
,”
J. Power Sources
,
296
, pp.
10
17
.
69.
Wang
,
W.-W.
,
Gu
,
Y.
,
Yan
,
H.
,
Li
,
S.
,
He
,
J.-W.
,
Xu
,
H.-Y.
,
Wu
,
Q.-H.
,
Yan
,
J.-W.
, and
Mao
,
B.-W.
,
2020
, “
Evaluating Solid-Electrolyte Interphases for Lithium and Lithium-Free Anodes From Nanoindentation Features
,”
Chem
,
6
(
10
), pp.
2728
2745
.
70.
Wang
,
M.
,
Huai
,
L.
,
Hu
,
G.
,
Yang
,
S.
,
Ren
,
F.
,
Wang
,
S.
,
Zhang
,
Z.
, et al
,
2018
, “
Effect of LiFSI Concentrations to Form Thickness- and Modulus-Controlled SEI Layers on Lithium Metal Anodes
,”
J. Phys. Chem. C
,
122
(
18
), pp.
9825
9834
.
71.
Hu
,
Z.
,
Wang
,
C.
,
Wang
,
C.
,
Chen
,
B.
,
Yang
,
C.
,
Dong
,
S.
, and
Cui
,
G.
,
2022
, “
Uncovering the Critical Impact of the Solid Electrolyte Interphase Structure on the Interfacial Stability
,”
InfoMat
,
4
(
3
), p.
e12249
.
72.
Xu
,
K.
,
2004
, “
Nonaqueous Liquid Electrolytes for Lithium-Based Rechargeable Batteries
,”
Chem. Rev.
,
104
(
10
), pp.
4303
4418
.
73.
Heine
,
J.
,
Hilbig
,
P.
,
Qi
,
X.
,
Niehoff
,
P.
,
Winter
,
M.
, and
Bieker
,
P.
,
2015
, “
Fluoroethylene Carbonate as Electrolyte Additive in Tetraethylene Glycol Dimethyl Ether Based Electrolytes for Application in Lithium Ion and Lithium Metal Batteries
,”
J. Electrochem. Soc.
,
162
(
6
), pp.
A1094
A1101
.
74.
Jung
,
R.
,
Metzger
,
M.
,
Haering
,
D.
,
Solchenbach
,
S.
,
Marino
,
C.
,
Tsiouvaras
,
N.
,
Stinner
,
C.
, and
Gasteiger
,
H. A.
,
2016
, “
Consumption of Fluoroethylene Carbonate (FEC) on Si-C Composite Electrodes for Li-Ion Batteries
,”
J. Electrochem. Soc.
,
163
(
8
), pp.
A1705
A1716
.
75.
Qian
,
J.
,
Henderson
,
W. A.
,
Xu
,
W.
,
Bhattacharya
,
P.
,
Engelhard
,
M.
,
Borodin
,
O.
, and
Zhang
,
J. G.
,
2015
, “
High Rate and Stable Cycling of Lithium Metal Anode
,”
Nat. Commun.
,
6
, p.
6362
.
76.
Zheng
,
J.
,
Lochala
,
J. A.
,
Kwok
,
A.
,
Deng
,
Z. D.
, and
Xiao
,
J.
,
2017
, “
Research Progress Towards Understanding the Unique Interfaces Between Concentrated Electrolytes and Electrodes for Energy Storage Applications
,”
,
4
(
8
), p.
1700032
.
77.
Chen
,
S.
,
Zheng
,
J.
,
Mei
,
D.
,
Han
,
K. S.
,
Engelhard
,
M. H.
,
Zhao
,
W.
,
Xu
,
W.
,
Liu
,
J.
, and
Zhang
,
J. G.
,
2018
, “
High-Voltage Lithium-Metal Batteries Enabled by Localized High-Concentration Electrolytes
,”
,
30
(
21
), p.
e1706102
.
78.
Li
,
N. W.
,
Yin
,
Y. X.
,
Yang
,
C. P.
, and
Guo
,
Y. G.
,
2016
, “
An Artificial Solid Electrolyte Interphase Layer for Stable Lithium Metal Anodes
,”
,
28
(
9
), pp.
1853
1858
.
79.
Wang
,
Q.
,
Wan
,
J.
,
Cao
,
X.
,
Wen
,
R.
,
Guo
,
Y.
,
Liu
,
W.
, and
Zhou
,
H.
,
2021
, “
Organophosphorus Hybrid Solid Electrolyte Interphase Layer Based on LixPO4 Enables Uniform Lithium Deposition for High-Performance Lithium Metal Batteries
,”
,
32
(
2
), p.
2107923
.
80.
Zhang
,
C.
,
Lan
,
Q.
,
Liu
,
Y.
,
Wu
,
J.
,
Shao
,
H.
,
Zhan
,
H.
, and
Yang
,
Y.
,
2019
, “
A Dual-Layered Artificial Solid Electrolyte Interphase Formed by Controlled Electrochemical Reduction of LiTFSI/DME-LiNO3 for Dendrite-Free Lithium Metal Anode
,”
Electrochim. Acta
,
306
, pp.
407
419
.
81.
Xu
,
R.
,
Cheng
,
X.-B.
,
Yan
,
C.
,
Zhang
,
X.-Q.
,
Xiao
,
Y.
,
Zhao
,
C.-Z.
,
Huang
,
J.-Q.
, and
Zhang
,
Q.
,
2019
, “
Artificial Interphases for Highly Stable Lithium Metal Anode
,”
Matter
,
1
(
2
), pp.
317
344
.
82.
Kwak
,
W.-J.
,
Park
,
J.
,
Nguyen
,
T. T.
,
Kim
,
H.
,
Byon
,
H. R.
,
Jang
,
M.
, and
Sun
,
Y.-K.
,
2019
, “
A Dendrite- and Oxygen-Proof Protective Layer for Lithium Metal in Lithium–Oxygen Batteries
,”
J. Mater. Chem. A
,
7
(
8
), pp.
3857
3862
.
83.
Kang
,
I. S.
,
Lee
,
Y.-S.
, and
Kim
,
D.-W.
,
2013
, “
Improved Cycling Stability of Lithium Electrodes in Rechargeable Lithium Batteries
,”
J. Electrochem. Soc.
,
161
(
1
), pp.
A53
A57
.
84.
Ma
,
G.
,
Wen
,
Z.
,
Wang
,
Q.
,
Shen
,
C.
,
Jin
,
J.
, and
Wu
,
X.
,
2014
, “
Enhanced Cycle Performance of a Li–S Battery Based on a Protected Lithium Anode
,”
J. Mater. Chem. A
,
2
(
45
), pp.
19355
19359
.
85.
Dong
,
H.
,
Xiao
,
X.
,
Jin
,
C.
,
Wang
,
X.
,
Tang
,
P.
,
Wang
,
C.
,
Yin
,
Y.
,
Wang
,
D.
,
Yang
,
S.
, and
Wu
,
C.
,
2019
, “
High Lithium-Ion Conductivity Polymer Film to Suppress Dendrites in Li Metal Batteries
,”
J. Power Sources
,
423
, pp.
72
79
.
86.
Zheng
,
G.
,
Lee
,
S. W.
,
Liang
,
Z.
,
Lee
,
H. W.
,
Yan
,
K.
,
Yao
,
H.
,
Wang
,
H.
,
Li
,
W.
,
Chu
,
S.
, and
Cui
,
Y.
,
2014
, “
Interconnected Hollow Carbon Nanospheres for Stable Lithium Metal Anodes
,”
Nat. Nanotechnol.
,
9
(
8
), pp.
618
623
.
87.
Gao
,
Y.
,
Zhao
,
Y.
,
Li
,
Y. C.
,
Huang
,
Q.
,
Mallouk
,
T. E.
, and
Wang
,
D.
,
2017
, “
Interfacial Chemistry Regulation Via a Skin-Grafting Strategy Enables High-Performance Lithium-Metal Batteries
,”
J. Am. Chem. Soc.
,
139
(
43
), pp.
15288
15291
.
88.
Hou
,
Z.
,
Zhang
,
J.
,
Wang
,
W.
,
Chen
,
Q.
,
Li
,
B.
, and
Li
,
C.
,
2020
, “
Towards High-Performance Lithium Metal Anodes Via the Modification of Solid Electrolyte Interphases
,”
J. Energy Chem.
,
45
, pp.
7
17
.
89.
Kang
,
D.
,
Xiao
,
M.
, and
Lemmon
,
J. P.
,
2020
, “
Artificial Solid-Electrolyte Interphase for Lithium Metal Batteries
,”
Batteries Supercaps
,
4
(
3
), pp.
445
455
.
90.
Liu
,
W.
,
Guo
,
R.
,
Zhan
,
B.
,
Shi
,
B.
,
Li
,
Y.
,
Pei
,
H.
,
Wang
,
Y.
,
Shi
,
W.
,
Fu
,
Z.
, and
Xie
,
J.
,
2018
, “
Artificial Solid Electrolyte Interphase Layer for Lithium Metal Anode in High-Energy Lithium Secondary Pouch Cells
,”
ACS Appl. Energy Mater.
,
1
(
4
), pp.
1674
1679
.
91.
Wu
,
P.
,
Dong
,
M.
,
Tan
,
J.
,
Kang
,
D. A.
, and
Yu
,
C.
,
2021
, “
Revamping Lithium-Sulfur Batteries for High Cell-Level Energy Density by Synergistic Utilization of Polysulfide Additives and Artificial Solid-Electrolyte Interphase Layers
,”
,
33
(
48
), p.
e2104246
.
92.
Yurkiv
,
V.
,
Foroozan
,
T.
,
Ramasubramanian
,
A.
,
Shahbazian-Yassar
,
R.
, and
Mashayek
,
F.
,
2018
, “
The Influence of Stress Field on Li Electrodeposition in Li-Metal Battery
,”
MRS Commun.
,
8
(
3
), pp.
1285
1291
.
93.
Li
,
J.
,
Dudney
,
N. J.
,
Nanda
,
J.
, and
Liang
,
C.
,
2014
, “
Artificial Solid Electrolyte Interphase to Address the Electrochemical Degradation of Silicon Electrodes
,”
ACS Appl. Mater. Interfaces
,
6
(
13
), pp.
10083
10088
.
94.
Zhao
,
Q.
,
Stalin
,
S.
, and
Archer
,
L. A.
,
2021
, “
Stabilizing Metal Battery Anodes Through the Design of Solid Electrolyte Interphases
,”
Joule
,
5
(
5
), pp.
1119
1142
.