All-solid-state lithium-ion batteries (ASSLIBs), with their exceptional attributes, have captured the attention of researchers. They offer a viable solution to the inherent flaws of traditional lithium-ion batteries. The crux of an ASSLB lies in its solid-state electrolyte (SSE) which shows higher stability and safety compared to liquid electrolyte. Additionally, it holds the promise of being compatible with Li metal anode, thereby realizing higher capacity. Inorganic SSEs have undergone tremendous developments in the last few decades; however, their practical applications still face difficulties such as the electrode–electrolyte interface, air stability, and so on. The structural composition of inorganic electrolytes is inherently linked to the advantages and difficulties they present. This article provides a comprehensive explanation of the development, structure, and Li-ion transport mechanism of representative inorganic SSEs. Moreover, corresponding difficulties such as interface issues and air stability as well as possible solutions are also discussed. Human civilization is heavily reliant on energy, but traditional fossil fuels such as coal, oil, and gas have a significant impact on the environment due to the release of carbon dioxide into the atmosphere, which is one of the primary contributors to global warming and poses a significant threat to the planet. Consequently, new energy sources such as wind and solar are gaining increasing importance as alternatives to traditional fossil fuels. However, these energy sources are variable in both space and time, necessitating the development of high-efficiency energy storage systems . Among t−1, while many solid-state batteries using Li metal anode in the laboratory have an initial discharge capacity of more than 100 mAh·g−1. However, the application of Li metal anode has been hampered by the growth of Li dendrites, unstable Li/electrolyte interfaces, and Li pulverization during battery cycling, which limits the capacity of conventional Li-ion batteries  skeleton is not suitable for Li+ migration. When the structure of the ionic conductor matches the size of migrating ions, the ionic conductor can achieve the maximum diffusion coefficient and the lowest activation energy [2.5Sr0.75Zr1.25(PO4)3 with impressive proton conductivity of 0.178S·cm−1 at 550 °C [2(PO4)3  and found that LiTi2(PO4)3 was the best NASICON-type electrolyte for Li-ion diffusion with the most suitable lattice size, the highest ionic conductivity, and the smallest activation energy for Li+ bulk migration [4+ in the structure exhibited poor stability especially after being in contact with metallic Li, another competitive candidate is LiGe2(PO4)3. The ionic conductivity of LiTi2(PO4)3 and LiGe2(PO4)3 solid electrolytes is significantly better than LiZr2(PO4)3. In addition to changing the bottleneck, another strategy to increase the conductivity of Li+ is the partial substitution of M4+ that increases the mobile lithium concentration and mobility [3+ to partially replace Ti4+ and prepared a series of Li1+xMxTi2-x(PO4)3, and the sample with M = Al or Sc showed high ionic conductivity of 7 × 10−4 S·cm−1 [4+ increases mobile lithium concentration and sinterability, thereby improving conductivity. Substitution of trivalent atoms not only increases the density of Li ions in the material but also induces additional interstitial migration with low activation energy, thereby affecting the mobility of Li ions [−3 S·cm−1 at room temperature [Casciola et al. studied the Lit 300 °C , which in energy . Lu et at 550 °C . Aono etigration . Neverthmobility . Aono et4 S·cm−1 . Trivale Li ions . The parperature , therefo1+xAlxTi2-x(PO4)3 (LATP) was reported by J. Fu in 1997, which exhibited an extremely high conductivity of 1.3 × 10−3 S·cm−1 at room temperature [1.3Al0.3Ti1.7(PO4)3 electrolyte sheet by a cold sintering process (CSP), and the wettability of the electrolyte sheet was improved. The grain boundary was densified by the method, with the total conductivity of the prepared LATP sheet improved from 3 × 10−6 S·cm−1 to 3 × 10−4 S·cm−1 [1.4Al0.4Ti1.6(PO4)3, which showed an ionic conductivity of 1.12 × 10−3 S·cm−1 and an activation energy of 0.25 eV at 25 °C [Liperature . Hamao e4 S·cm−1 . Xu et aat 25 °C . 4+, which readily reduces to Ti3+ when in contact with the Li anode. Although Li1+xAlxTi2-x(PO4)3 (LAGP) does not show as high conductivity as LATP, the high stability of LAGP makes it a popular material as an electrolyte. In 1997, Fu et al. reported Li1.5Al0.5Ge1.5(PO4)3 glass-ceramics with a conductivity over 10−4 S·cm−1 at room temperature [1.5Al0.5Ge1.5(PO4)3 by glass sintering with concurrent crystallization [−4 S·cm−1 and an electrochemical stability window around 3.5 V. Crystallization treatment is beneficial for obtaining glass-ceramics with higher electrical conductivity by controlling microstructures [−4 S·cm−1 [2O as a secondary phase in LAGP, which acted as a nucleating agent to significantly promote the crystallization of the as-prepared glasses during heat treatment, leading to an improvement in the connection between the glass-ceramic grains and resulting in a dense microstructure with a uniform grain size [2 sample exhibits an ionic conductivity of 7.25 × 10−4 S·cm−1 and a stability window as high as 6 V. Nikodimos et al. enhanced the ionic conductivity of LGP by substituting 25% of Ge4+ ions in the LGP structure with Al3+ and Sc3+ ions, which leads to more Li+ in the vacant sites. The sample with the composition of Li1.5Al0.33Sc0.17Ge1.5(PO4)3 shows the highest bulk Li+ conductivity of 5.826 × 10−3 S·cm−1 with Ea of 0.279 eV [The electrochemical window of the LATP-type material is limited by the low stability of Tiperature . Zallocclization . Their sructures . Zhu et 4 S·cm−1 . Xu et aain size . The LAG0.279 eV .4+ is easily converted to Ti3+, which causes LATP to degrade when in contact with the Li anode and lower ionic conductivity at the electrolyte–electrode interface. By replacing Ti4+ with Ge4+, LAGP is more stable and exhibits a wide electrochemical stability window of 0.85–7 V vs. Li/LATP has low stability toward metallic Li because Tis. Li/Li+. Althougs. Li/Li+. When cos. Li/Li+. The mets. Li/Li+. Doping 4+, Hf4+) appear to be suppressing the Ti4+ reduction, whereas large divalent cations  appear to be promoting it even at low concentrations. Doping of Mg2+ showed improvement in the Ti4+ reduction tolerance; however, it demonstrated a detrimental effect on the conductivity after contact with Li metal [2+ in LATP and found that the Mg2+ did not bleed heavily into the adjacent crystalline grain domains, which made it a suitable dopant for interfacial engineering. The small amount of Mg2+ ions that dope into the grain shows preferential substitution in the Ti/Al host framework, which preserves the functionality of the 3D Li network in LATP. Reduced Ti4+ content may, by polaron hopping, result in significantly lower residual electrical conductivity and shield the electrolyte from degradation [2F2 to modify the defect sites of the LATP pellet and impede the interfacial reactions between the LATP electrolyte and Li anode [Mashekova et al. studied the effect of tetravalent and divalent cation dopants  of LATP on the Li-ion conduction and Ti reduction during interaction with lithium metal. The conductivity is not significantly impacted by a small ratio of doping, but its redox property is altered. Similar-sized tetravalent cations 30 electrolyte performed over 600 h at 0.1 mA·cm−2, and the modified LATP exhibited electrochemical stability up to 5 V [2O3 coated by atomic layer deposition technique as 2O3-coated LATP exhibited a stable cycling behavior with smaller voltage hysteresis for 600 h, as well as small resistance. More importantly, the lithium penetration into the LATP bulk and Ti4+ reduction were significantly limited [2 coating layer as an ASEI on LATP. The MoS2 layer not only effectively inhibits the decomposition of LATP, but also forms in situ a Li2S and Mo metal conversion layer during cycling, which can improve the interfacial charge transfer kinetics and decrease the charge transfer resistance [−2 from ∼30 h without protection [2OHBr ionic conductor as a protective layer between LAGP and Li anode [2OHBr layer not only facilitates the migration of Li ions but also effectively avoids the reduction of LAGP by Li metal. As the most commonly used method to improve battery cycle performance, interface engineering is also applicable. As shown in p to 5 V . Liu et  limited . Huang esistance . The assotection . The Cr Li anode . The Li2+ ion moves with Co3+ to LATP, which led to high Li chemical potential around the LATP/LCO interfaces and the dynamical Li-ion depletion upon charging [2(PO4)3 (M = Fe and others), which remains as a thermodynamically stable phase at high temperatures in contact with the LATP [2+ migrates from the nickel-rich cathodes to the NCM/LATP interface. Subsequently, Ni2+ reacts with Ti4+ in LATP, resulting in the decomposition of LATP. Wang et al. constructed a perovskite La4NiLiO8 buffer layer between NCM and LATP to reduce the content of Ni2+ on the surface, which enhanced the interfacial stability [During the charge and discharge process, ion exchange may happen between the electrolyte and cathode. Between LATP and LCO, Co-Ti exchange was found as Licharging . As a recharging . In addithe LATP . When LAtability . 16-2xDx(TO4)4, where D = Mg2+ or Zn2+, T = Si4+ or Ge4+, and 0 < x < 4 [14Zn(GeO4)4 exhibits ultrahigh conductivity 0.125 S·cm−1 at 300 °C, thus the structure was named Li superionic conductor (LISICON). The structure has the space group of Pnma and is similar to γ-Li3PO4 Each O2− is bonded to four network cations  to form a rigid three-dimensional network of Li11Zn(GeO4)4. The three remaining Li+ ions are weakly bonded with O2− and are free to diffuse between the interstitial positions in the network. The average diameter of the bottlenecks between these positions is larger than twice the sum of the ionic radii of Li+ and O2−, satisfying the geometrical requirement for rapid Li+ transport. Ion transport in the structure is two-dimensional because the connected sites of mobile Li+ lie in the same plane. However, the conductivity of Li14Zn(GeO4)4 is only ~10−6 S·cm−1 at room temperature [In 1978, Hong et al. were the first to synthesize a series of Li < x < 4 . Among tperature . + mobility appears to be favored by the introduction of interstitial Li+ ions in γ-tetrahedral structures. Hu et al. prepared Li4-xSi1-xPxO4 (0 < x < 1), the solid solutions formed between Li4SiO4 and Li3PO4. The samples with x = 0.5 and 0.6 were found to have the highest conductivity [4SiO4 was added to Li3PO4, and the number of Li+ ions occupying interstitial positions increased, thereby improving the ionic conductivity. Kuwano et al. synthetized the Li-ion conductor of the Li4GeO4-Li3VO4 system which shows higher ionic conductivity. The total conductivity of Li3.6Ge0.6V0.4O4 reached 4 × 10−5 S·cm−1 at 18 °C which is much higher than the previous LISICON structure [+ ions which are created during solid solution formation. In the γ-Li3VO4 cell, V5+ is partially substituted by Li+ and Ge4+ and additional Li+ ions occupy the interstice sites to form interstice Li+ ions in the Li3+xGexV1-xO4 crystal. The ionic conductivity is enhanced by the abundant Li+ ions in the interstice. Additionally, Li3+xGexV1-xO4 exhibits good thermal stability and is stable when in contact with the atmosphere of CO2. Deng et al. simulated and prepared a series of Si-substituted Li4±xSi1-xXxO4 . From low to high temperatures, three Li-ion diffusion mechanisms were identified: local oscillation, isolated hopping, and superionic motion. The substitution of P, Al, or Ge on Si sites caused polyanion mixing, which lower the temperature at which the transition to a superionic state and improved the room Li-ion conductivity by several orders of magnitude [3.75±y(Ge0.75P0.25)1-xMxO4  and Li3.75-y(Ge0.75P0.25)1-xVxO4 (x = 0–0.5) to assess the effect of cation co-doping in LISICON structure, and Li3.53(Ge0.75P0.25)0.7V0.3O4 showed the highest ionic conductivity of 5.1 × 10−5 S·cm−1 at 25 °C. V5+co-doping lowers the energy barrier for Li+ diffusion in the lattice, which further increases ionic conductivity [For the LISICON structure, high Liuctivity . Si parttructure . High coagnitude . Zhao etuctivity .− for O2−. As the radius of Cl− is bigger than that of O2−, the substation of Cl− for O2− increases the lattice constants thus enlarging the bottleneck size, lowering the activation energy, and enhancing the ionic conductivity. However, the ionicity of Li-Cl is smaller than that of Li–O due to the lower electronegativity of Cl−. With increasing Cl concentration, the Li+ become tightly bonded to the Cl−, acting as an impediment to ion migration, increasing activation energy, and decreasing ionic conductivity. Therefore, appropriate Cl− doping can significantly improve the properties of LISICON-type conductors. It was reported that Li10.42Si1.5P1.5Cl0.08O11.92 and Li10.42Ge1.5P1.5Cl0.08O11.92 show the highest ionic conductivities of 1.03 × 10−5 S·cm−1 and 3.7×10−5 S·cm−1 at 27 °C, respectively, which is an order of magnitude higher than Li10.5Ge1.5P1.5O12. With the introduction of Cl−, the electrochemical stability with lithium metal up to 9 V vs. Li+/Li, which is one of the widest electrochemical windows for solid electrolytes [4GeO4 will exhibit the highest ionic conductivities at 100 °C among the traditional LISICONs if it can be synthesized [The effect of anionic substitution in the LISICON structures was also investigated. Song et al. reported a method for achieving high conductivity and great electrochemical stability by the substitution of Cltrolytes . Fujimurthesized . Overall2− in LISICON-type solid electrolytes with S2−. Past research has shown that the size of the S2− is near optimal for Li+ conduction in this structural framework [The substitution of cations in the LISICON structure does not significantly improve the electrochemical performance, while the substitution of anions greatly changes the performance. The ionic conductivity can be significantly improved by replacing Oramework . 2S-GeS2, Li2S-GeS2-ZnS, Li2S-GeS2-Ga2S3 systems, which showed the highest conductivity of 6.5 × 10−5·S·cm−1 at room temperature [+. Liu et al. synthesized Li2ZrS3 through a solid-state reaction method and obtained Li2ZrS3 had ionic conductivity of 7.3 × 10−6 S·cm−1 at 30°C. With Zn2+ doping and partially substituting Zr4+, the Li2.2Zn0.1Zr0.9S3 exhibited much higher ionic conductivity of 1.2 × 10−4 S·cm−1 [3+5xP1-xS4 ∞ chain. Along the one-dimensional chain, the Li(1)S4 tetrahedrons are connected by corners, and an interstitial tetrahedral site exists between two Li(2)S6 octahedrons and one Li(3)S4 tetrahedron [3PS4 with nanoporous structure, which exhibited a wide electrochemical window (5 V), superior chemical stability against lithium metal as well as an anomalous high ionic conductivity of 1.6 × 10−4 S·cm−1 at 25 °C, 3 orders higher than the intrinsic ionic conductivity of bulk β-Li3PS4 (8.93 × 10−7 S·cm−1). The nanoporous structure of β-Li3PS4 combines two important factors that improve ionic conductivity: (1) the stabilization of the high-conduction phase that develops at high temperatures and (2) the promotion of surface conduction by the high surface-to-bulk ratio of nanoporous β-Li3PS4 [Homma et al. reported three phase transitions (γ → β → α) of Lirahedron . Liu et β-Li3PS4 . As the β-Li3PS4 . 10GeP2S12 (LGPS) with ionic conductivity of 1.2 × 10−2·S·cm−1 at room temperature, which was much higher than the solid-state Li conductors ever reported in recent decades and comparable to organic liquid electrolytes [2/nmc. As shown in 6 octahedra and (Ge0.5P0.5)S4 tetrahedra connected by a shared edge to create one-dimensional (1D) chains, which form the framework of LPGS [+ ions have two migration modes: along the tunnel in the c-axis and in the ab plane. The migration of Li+ along the c-axis is more favorable than that on the ab plane, which leads to weak anisotropy of Li+ migration of LGPS [−3 S·cm−1, which is comparable to advanced solid electrolytes [+ ion diffusions in the LGPS by 7Li and 31P multiple solid-state NMR methods. + ion diffusion modes and the activation energy for diffusion in the tunnel along the c-axis is 0.16 eV, which is much lower than diffusion in the ab plane (Ea = 0.26eV) [+ ions may be the origin of the collective migration process that results in the high conductivity along the c axis. As +(1) hops to fill the adjacent vacancy, a large Coulombic repulsion against Li+(2) is generated, which prompts the next Li+ to jump to the latter position. In the synchronized action, the Li+ ions replace one another in a string-like fashion and this collective ionic migration exhibits a fairly low activation energy barrier [4GeS4−yLi3PS4 with different x/y were studied (4/1 ≥ x/y ≥ 1/2). The room-temperature Li ion conductivity increases and subsequently declines as the Ge/P ratio increases, peaking at x/y = 1/1.2. The system with higher Ge content forms an ordered LGPS structure with slightly higher activation energy at room temperature [In 2011, Kamaya et al. reported a new thio-LISICON-type Litrolytes . LGPS ha of LPGS . In LGPS of LGPS . Althougtrolytes . Liang e 0.26eV) . The str barrier . Additioperature .3.25Ge0.25P0.75S4 through a low-temperature solution method, and the LGPS film exhibited a lithium-ion conductivity of 1.82 × 10−4 S·cm−1. Though the conductivity of the film is lower than bulk LGPS due to the film’s stoichiometry deviation and poor crystalline nature, it has a lithium-ion transference number of >0.999, which could be applicable in the thin film all-solid-state lithium-ion batteries [Research on the application of LGPS has been put forward. Wang et al. synthesized Liatteries . Dawson atteries . 2− in LISICON-type solid electrolytes being replaced by S2−, thio-LISICON structure electrolytes show great improvement. Despite the necessity to address problems such as the high price of Ge and poor stability in damp air, LGPS is thought to have excellent application prospects since it exhibits an ionic conductivity comparable to organic liquid electrolytes. Generally speaking, with O3P, Li2S, and Li-Ge alloy, which increase the interfacial resistance [3.25Ge0.25P0.75S4) spontaneously form a breathing interface with the application of electric current, which effectively makes close contact at the electrolyte/electrode boundary and promotes fast charge transfer at the solid interface. For the Li anode, the SEI phase between the Li/thio-LISICON interface gradually grows during cycling and ultimately destroys the cell. Another study about Li-Al alloy anode also verified the excellent compatibility of Li-Al alloy and LGPS. Li0.8Al|LGPS|Li0.8Al could operate stably for more than 2500 h at 0.5 mA·cm−2 [2O)n-Li) and inorganic nanoparticle salts  as interphase to protect LGPS from decomposition [2PO4 protective layer on the Li anode to address intrinsic chemical stability problems of LGPS toward Li metal [2|LGPS|Li ASSLiB showed a reversible discharge capacity of 131.1 mAh·g−1 at the initial cycle and 113.7 mAh·g−1 at the 500th cycle under 0.1 C with a retention of 86.7%.Sulfide electrolytes frequently do not have good stability with Li metal, and unstable interfaces can form. Take LGPS as an example, the intrinsic stability window of LGPS is very limited, from 1.7 to 2.1 V . It was sistance . More imsistance . Thus, Lsistance . Li-Al a mA·cm−2 . Gao et position . Figure Li metal . The sta−2 without lithium dendrite penetration or short circuit [Another crucial concern is the formation of lithium dendrites in solid-state batteries when the current density exceeds a threshold amount . Introdu circuit . Additio circuit .2. After cycling, side products such as sulfides are observed, accompanied by an increase in interfacial resistance and capacity fading. Three possible situations between LGPS and LiCoO2 during long-term cycling are proposed as +/Li < E < 4.2 V vs. Li+/Li), LGPS decomposed with the formation of polysulfide species; at high potential (E ≥ 4.2 V vs. Li+/Li), the high potential triggers the structural instability and oxygen release at the cathode, which results in the formation of phosphate and sulfate species and more severe degradation [Side reactions at the cathode during cycling are the main causes of electrolyte–cathode interface degradation. Zhang et al. studied the degradation of the interface between LGPS and LiCoOradation . In addiradation .2S [2S gas, which leads to the collapse of structure and a decrease in conductivity [4+, As5+, Sb5+ are inclined to bind tightly with the soft base S2− and impede hydrolysis by H2O [5+ and hard base O2−, the introduction of O2− into the structure shows significant improvement of air stability toward moisture in the air.Although the ionic conductivity of thio-LISICON is significantly improved, thio-LISICON is not stable when exposed to air and moisture. The practical applications of thio-LISICON without modification are still limited. The P-S bond in the structure is unstable and can be easily replaced by other atoms to form a stable building block. As 2S . Therefouctivity . The moss by H2O . Similar4–xSbxSn1–xS4 with 0 ≤ x ≤ 0.5 [3.8Sb0.2Sn0.8S4 shows the highest ionic conductivity of 3.5 × 10−4 S·cm−1 among the system and high stability in moist air due to the high affinity of Sn-S and Sb-S. Liang et al. doped Sb5+ into LGPS as Sb5+ can form stronger covalent bonds with S2− than P5+ [10Ge(P1–xSbx)2S12 ionic conductors exhibit high air stability and have a high conductivity of 12.1~15.7 × 10−3 S·cm−1 even after being exposed to humid environment. Zhang et al. reported the solid solutions of Li x ≤ 0.5 . Li3.8Sbthan P5+ . The Li12− for S2− in the structure can not only maintain the high conductivity of the sulfide prototype, but also improve the stability. Tsukasaki et al. replaced Li2S with Li2O in the 75Li2S·25P2S5 glass. The (75-x)Li2S·25P2S5·xLi2O could still maintain an ionic conductivity of higher than 10−4 S·cm−1 and better stability to air. For the samples with x ≥ 15, a good balance of thermal stability, wet resistance, and ionic conductivity can be expected [9+δP3+δ′S12–kOk series of solid solution phases and proposed that the solid solution range was 0 < k ≤ 3.6 based on the lattice parameter variation [−3 S·cm−1 as well as good stability in the air atmosphere. Liu et al. prepared a series of Li10Sn0.95P2S11.9−xOx (0 ≤ x ≤ 1) by the solid-phase sintering method and Li10Sn0.95P2S11.4O0.5 shows high conductivity of 3.96 × 10−3 S·cm−1 with a negligible decrease after exposed in the air [5+ and O2− substituted Li10SnP2S12 with high air stability due to the soft acid Sb5+ and hard base O2− dual substitution [10SnP1.84Sb0.16S11.6O0.4 electrolyte displays a broader electrochemical window of 1.4–5.0 V vs. Li+/Li and ionic conductivity of 2.58 × 10−3 S·cm−1.The partial substitution of Oexpected . Xu et aariation . The sam the air . Gao et titution . The Li12S gas, resulting in a decrease in conductivity, Weng et al. reported a protective layer deliberately forming on the LGPS surface by controlling humidity [4P2S6, GeS2, and Li2HPO3 rapidly forms in dozens of seconds. The forming layer, which is ionically conductive and electronically insulating, can significantly prevent the severe interface reaction between LGPS and Li anode during electrochemical cycling. Both the symmetrical cell and full cell with LGPS exposed for 40 s exhibit good cycle performance. Although it is widely recognized that the water in the air can lead to severe decomposition of LGPS and release of Hhumidity . In thisIn this article, four typical inorganic solid-state electrolytes are introduced. Due to the difference in structure and composition, each solid electrolyte has its advantages and difficulties. The conductivities of mentioned electrolytes are summarized in −3 S·cm−1 through doping strategy. Al-LLZO, Ga-LLZO, and Ta-LLZO exhibit high conductivity and have been extensively studied. Most garnets have relatively good chemical stability to Li; however, Ga-LLZO is prone to undesirable side reactions with Li. In addition, Garnet-type electrolyte is not stable in contact with moisture and CO2 in the ambient atmosphere. Therefore, the use of Garnet-type electrolytes in a battery needs modification to improve stability.Garnet-type electrolytes with prototype of LLZO can achieve high conductivity of ~10−4 or 10−5 S·cm−1. LLTO has poor stability towards Li and Ti-free perovskite-type electrolytes can solve this problem to some extent. Perovskite-type electrolyte exhibits high compatibility with commonly used cathode materials and can be used as coating layer on cathodes.The application of perovskite-type electrolytes is limited due to low total conductivity led by high grain boundary resistance. The conductivity of perovskite electrolytes is only on the order of 10As the most representative NASICON-type electrolytes, LATP and LAGP show high ionic conductivity and moisture/air stability; however, they have low stability towards Li anode. The commercialization of LAGP is limited by the high cost of Ge-containing raw materials.−5 S·cm−1. Thio-LISICON-type electrolytes show much higher ionic conductivity with O2− replaced by S2−. Among them, LGPS has the highest conductivity of ~10−2 at room temperature but low air stability due to its reaction with moisture. After partial substitution based on HSAB theory, the electrolyte still has a sufficiently high conductivity of ~10−3 with higher stability.LISICON-type electrolytes exhibit good chemical and electrochemical stability but low conductivity of ~10−3~10−2 S·cm−1. Some of the solid-state electrolytes show comparable ionic conductivity and are considered to be promising candidates for all-solid-state Li-ion battery.Commercial liquid electrolytes usually use DMC, EC, or PC as mixed solvents, LiPF6 as ion-conductive salt, and have the conductivity of the order of 10For all-solid-state Li-ion batteries with Li as anode, the formation of dendrites is a problem that cannot be ignored. Solid-state electrolytes with high density can resist penetration by Li dendrites to a certain extent; however, the risk of short circuit persists after a long-time cycle. Therefore, it is necessary to make the current even distribution and regulate the deposition of Li by interface modification. The performances of mentioned cell with modified electrolytes are summarized in −1. CALT recently reported a Qilin battery with an energy density of 255 Wh·kg−1, which can be converted to a capacity of 63.75 mAh·g−1 at 4 V. As can be seen from −1 with an operating voltage of about 4 V, which is much higher than commercial lithium-ion batteries with liquid electrolytes. However, the cycle perforsmance of some solid-state batteries is poor, which means that there is still a certain distance from commercialization standards for many batteries in the laboratory. Although some Li-ion batteries with solid-state electrolytes in the laboratory show good performance, it is challenging to commercialize them because of their high cost.The commercial lithium-ion secondary batteries produced by Panasonic have capacity of about 60~70 mAh·gThis review concludes the research on four types of representative inorganic SSEs: garnet, perovskite, NASICON, and thio-/LISICON. The crystal structures of these SSEs, Li-ion transport mechanism, and ionic conductivity improvement techniques are all briefly discussed. Interfaces between solid-state electrolytes and anode/cathode, air stability of solid-state electrolytes, and associated improvement strategies are also addressed.Compared to liquid electrolytes, SSEs have higher stability and safety as well as the prospect of being compatible with Li metal anode to achieve higher capacity. However, the rigid nature of solid-state electrolytes and poor contact of solid–solid interface lead to high interface resistance and non-uniform current distribution at the interface. Furthermore, the battery’s ability to retain high capacity and long-term processing has always been hampered by issues with the stability of the electrolyte and electrodes at their interface and stability in the air. Although substantial achievements have been made so far, there is still a long way to go before the successful industrialization and commercialization of all-solid-state lithium-ion batteries. Based on the aforementioned research, the following elements of future solid-state electrolytes need to be improved.(1) Improving the ionic conductivity of solid-state electrolytes: Most solid electrolytes still lag far behind organic liquid electrolytes in ionic conductivity. On the one hand, the material composition, structure, and doping strategy of ionic conductors directly determine the ionic conductivity of the material. On the other hand, the preparation process affects the properties of the obtained product. It is important to optimize material composition and preparation process in order to improve ionic conductivity.(2) Improving the stability of solid-state electrolytes toward the Li metal anode: Compatibility with high-capacity Li electrodes is one of the advantages of solid-state electrolytes, though the use of lithium metal electrodes is subject to certain limitations. The inclusion of an appropriately modified layer can minimize the growth of Li dendrites while also significantly enhancing the stability and contact of the interface. Additionally, the use of a lithium alloy anode can significantly lower the production of Li dendrites, and the stability of electrolytes can be increased by using the right doping and manufacturing techniques. (3) Improving the stability of solid-state electrolytes toward cathodes: The reactions at the interface between cathodes and electrolytes are very complex. At the interface, the cathodes and electrolytes break down, which might create a layer of depleted lithium and cause the interface to degrade. Exploring electrolytes with higher oxidation stability is one solution. In addition, constructing a suitable artificial solid-electrolyte interface layer also contributes to the improvement of stability.2O and/or CO2 in the air. Sulfide electrolytes are particularly poorly stable to moisture in the air. Doping modification is the main method to improve the stability to the air, and attention should also be paid to the maintenance of electrolytes’ properties.(4) Improving the air stability of solid-state electrolytes: When electrolytes are exposed to air for a long period, they may react with the H(5) Reducing costs and optimizing processes: Through unique synthesis techniques and interface alterations, solid-state electrolytes have been improved in numerous previous studies, making them more appropriate for all-solid-state battery operations over a lengthy period. However, some methods greatly increase the cost and complexity of production, which is not suitable for practical industrialization and commercialization.Numerous types of electrolyte materials have been proposed as a result of the growing depth of the study on solid electrolytes. However, the intrinsic properties of these materials cannot fully meet all the requirements of solid-state batteries. The stability of the electrode–electrolyte interface is the most difficult issue to solve, and it will also be the subject of a future study. On the one hand, it is required to enhance the intrinsic qualities of electrolyte materials to make the procedure easier and the cost of modification lower. On the other hand, the practical application must investigate optimization techniques that can be used in commercial and industrial production.


Nonetheless, the thermoelectric performance is poor and remains subpar. Herein, the electrical and thermal transport properties of Cu3SbSe4 were synergistically optimized by S alloying. Firstly, S alloying widened the band gap, effectively alleviating the bipolar effect. Additionally, the substitution of S in the lattice significantly increased the carrier effective mass, leading to a large Seebeck coefficient of ~730 μVK−1. Moreover, S alloying yielded point defect and Umklapp scattering to significantly depress the lattice thermal conductivity, and thus brought about an ultralow κlat ~0.50 Wm−1K−1 at 673 K in the solid solution. Consequently, multiple effects induced by S alloying enhanced the thermoelectric performance of the Cu3SbSe4-Cu3SbS4 solid solution, resulting in a maximum ZT value of ~0.72 at 673 K for the Cu3SbSe2.8S1.2 sample, which was ~44% higher than that of pristine Cu3SbSe4. This work offers direction on improving the comprehensive TE in solid solutions via elemental alloying.Cu S2σT/κ, where S, σ, T, and κ stand for the Seebeck coefficient, electrical conductivity, absolute temperature in Kelvin, and total thermal conductivity (comprising lattice part κlat and electronic part κele), respectively [S2σ) and/or lower κ [S, σ, and κele in the given TE material due to their strong coupling effects [κlat stands as the sole independently regulated TE parameter, leading to extensive research over the last two decades [Thermoelectric (TE) technology has the capability to directly and reversibly convert heat into electricity, making it a promising source of clean energy. It plays a significant role in addressing the challenges posed by the energy and environmental crises ,2,3. Numectively ,13. Actu lower κ ,18,19,20 effects ,21. Neve decades ,17,19.3SbSe4 compound is a p-type semiconductor, featuring a narrow band gap of ~0.29 eV [κ and low σ, stemming from low carrier concentration and mobility, present challenges that hinder its practical use. Extensive efforts have been implemented to enhance the TE performance of Cu3SbSe4, including elemental doping [n), S, or κlat, and thus leading to an appealing figure of merit. Although high n can enhance σ, it has a negative impact on S and result in an increase in κele. The TE performance of Cu3SbSe4 falls significantly short of that of Cu-based chalcogenides due to these two inherent issues. On one side, the narrow energy band gap of ~0.29 eV leads to bipolar diffusion, causing deterioration in electrical properties [κlat) inherently arises from its composition comprising lightweight elements and a diamond-like structure [Copper-based chalcogenides have garnered significant attention because of their relatively favorable electrical transport and low thermal transport properties ,23,24,25~0.29 eV . More im~0.29 eV ,32. Howel doping ,35,36,37l doping ,39,40, al doping ,42. Thesoperties . On the κlat and thereby enhancing the TE performance. For example, Skoug et al. demonstrated that the substitution of Ge on Sn sites can lead to the formation of Cu2Snx1−GexSe3 solid solutions, synergically optimizing the TE properties [max value of ~0.42 was obtained in the Cu2Ge(Sx1−Sex)3 system via Se alloying [2Ge(Sex1−Tex)3 by incorporating Te on the Se site, resulting in a ZTmax of ~0.55, which was 62% higher than that of the matrix [3Sb(Sex1−Sx)4 solid solution is an effectively strategy for enhancing the thermoelectric performance of the Cu3SbSe4 compound via S alloying. Moreover, the development of TE materials with more cost-efficient constituent elements is of significant importance for large-scale practical applications.The formation of a solid solution via elemental alloying is an effective strategy for depressing the operties . Jacob ealloying . Wang ete matrix . The afo3Sb(Sex1−Sx)4 solid solutions with x covering the whole range from 0 to 1. The results demonstrate that the Cu3SbSe4-Cu3SbS4 solid solutions exhibit an extremely high Seebeck coefficient and ultralow thermal conductivity. Firstly, S alloying can widen the band gap, alleviating the bipolar effect. Additionally, S substitution in the lattice can significantly increase the carrier effective mass, leading to a remarkably high Seebeck coefficient of ~730 μVK−1. Moreover, the κlat can be significantly depressed owing to point defect scattering and Umklapp scattering, thus obtaining a minimum κlat of ~0.50 Wm−1K−1. Consequently, the multiple effects of S alloying boost the TE performance of the Cu3SbSe4-Cu3SbS4 solid solution, and a maximum ZT value of ~0.72 at 673 K is obtained for the Cu3SbSe2.8S1.2 sample.Herein, we present the synthesis and thermoelectric characterization of the Cu3Sb(Sex1−Sx)4 solid solutions with varying S content  were synthesized by vacuum melting and plasma-activated sintering . Concretely, the synthesis was divided into two steps. The first step was to synthesize the primary powders. Firstly, the starting materials, consisting of high-purity components  corresponding to the nominal composition of Cu3Sb(Sex1−Sx)4 (x = 0–1), were carefully sealed in the quartz tube under high vacuum conditions (<10−3 Pa). Afterwards, the sealed tubes were incrementally heated to 1173 K with a controlled rate of 20 K/h and maintained at 1173 K for a duration of 12 h. Following a holding period, the tubes were cooled down with a relatively low rate of 10 K/h until reaching 773 K, and finally the samples were quenched into water. Subsequently, the acquired quenched ingots underwent direct annealing at 573 K for a period of 48 h to facilitate the uniformity of chemical compositions. After this step, the obtained ingots were finely pulverized using an agate mortar to produce uniform powders. The second step was to synthesize the target samples. The resultant powders were then introduced into a graphite die of Ø12.7 mm in diameter and treated using the PAS technique at 673 K for a duration of 5 min while applying an axial pressure of 50 MPa. In detail, the sintering temperature reached to 523 K after an activation time of 10 s under the activation voltage of 20 V and the activation current of 300 A, and then the current was manually adjusted to increase by a rate of 1.5 K/s to reach the desired sintering temperature of 673 K after 225 s; the temperature was then held for 300 s. Ultimately, the samples were furnace-cooled to room temperature.The Cu3Sb(Sex1−Sx)4 (x = 0–1) solid solutions were conducted using a Bruker D8 advance instrument, which was equipped with Cu Kα radiation (λ = 1.5418 Å). Lattice parameters were refined using the Rietveld method, employing the HighScore Plus computer program for analysis. The morphologies and compositions of the afore-mentioned solid solutions were performed by a Nova NanoSEM450 (FESEM) and a JEM-2010F (HRTEM), equipped with a detector of energy-dispersive X-ray spectroscopy (EDS).The X-ray diffraction (XRD) patterns for the Cuσ and S by the commercial measuring system  under a helium atmosphere, spanning a temperature range from room temperature to 673 K. Thermal conductivity was calculated using the equation of κ = DCpρ. Herein, the D, Cp, and ρ stand for the thermal diffusivity, specific heat, and density, respectively. The disks were used for simultaneously measuring D and Cp by utilizing a Laser Flash apparatus of Netzsch (LFA-457) under a static argon atmosphere. The ρ of the Cu3Sb(Sex1−Sx)4 (x = 0–1) solid solutions were conducted using Archimedes’ methods. The relative densities, in relation to the theoretical density of 5.86 g cm−3, have been provided in n (carrier concentration) and μ (carrier mobility) of the afore-mentioned solid solutions at 300 K were performed using the Hall effect system  according to the van der Pauw method under a magnetic field strength of 0.68 T.The as-sintered cylinders were processed into bars of 10 mm × 2 mm × 2 mm and disks of Ø12.7 mm × 2 mm. The bars were used for concurrently measuring 3Sb(Sex1−Sx)4 (x = 0–1) samples were performed by XRD. 3SbSe4, with blue, gray, and green atoms representing Cu, Sb, and Se, respectively. As displayed in x = 0) are fully indexed to the zinc-blende-based tetragonal structure (I-42m space group) of Cu3SbSe4 (JCPDS No. 85-0003) without any detectable impurities [x < 1), a continuous shift of the (112) diffraction peak towards higher angles can be seen (3Sb(Sex1−Sx)4 solid solutions. The shift in the diffraction peak can be ascribed to the smaller radius of S2− (1.84 Å) in comparison to Se2− (1.98 Å) [x = 1, the XRD peaks match the pattern of Cu3SbS4 (JCPDS No. 35-0581) [The crystal structures and phase compositions for the Cupurities . With in be seen c, demons(1.98 Å) . For x =35-0581) .3Sb(Sex1−Sx)4 (x = 0.3) samples based on the famatinite crystal structure are shown in pR, wpR, and expR) of Cu3Sb(Sex1−Sx)4 (x = 0–1) samples are listed in 3SbSe4-Cu3SbS4 solid solutions.The Rietveld refinement profiles of the Cu3Sb(Sex1−Sx)4 (x = 0.3) sample were characterized by a SEM equipped with an EDS detector (x = 0.3) indicated that they were isotropic materials. The nanopores (marked by the blue dotted circles) were observed on the fracture surface due to the Se/S volatilization of the synthesis process of the sample (3SbSe4-Cu3SbS4 (x = 0.3) solid solution.The morphologies and chemical compositions of the Cudetector . As prese sample a, which e sample c. Accorde sample d–h, the e sample , which d3Sb(Sex1−Sx)4 (x = 0.3) were further investigated at nanoscale using high-resolution TEM (HRTEM)  . The TEMsolution . As presectively . Additioectively .3Sb(Sex1−Sx)4 (x = 0–1) solid solutions, the charge transport properties were conducted. The temperature dependence of electrical conductivity (σ) of the Cu3Sb(Sex1−Sx)4 (x = 0–1) solid solutions is displayed in 3SbSe4 exhibited a monotonous increase in σ with rising temperature, demonstrating characteristic behavior of a non-degenerate semiconductor. For the x > 0.2 samples, the samples showed a transition from non-degenerate semiconductors to a partially degenerate regime [σ exhibited an initial decrease followed by an increase, with the minimum value occurring at ~473 K, indicating its association with bipolar conduction [σ of S alloying samples increased with the S contents until x = 0.3, after which it started to decrease with a higher S content. Notably, the σ improved from ~4.6 S/cm of pristine Cu3SbSe4 to ~42 S/cm of x = 0.3 solid solution at room temperature, arising from the augmented carrier concentration (x > 0.5) had lower σ compared to the pristine Cu3SbSe4, which was ascribed to the reduced n (carrier concentration) and diminished μ (carrier mobility). Furthermore, due to the intensified lattice vibration at elevated temperatures, the solid solutions exhibited lower σ than the pristine sample at high temperatures, indicating that the intensified lattice vibration in the solid solutions at elevated temperatures hindered the carrier migration [To explore the effects of S alloying on the TE properties of the Cue regime . The σ enduction ,52. The ntration . It is wigration ,53.S of the Cu3Sb(Sex1−Sx)4 (x = 0–1) samples. The p-type semiconductor behavior of solid solutions, characterized by dominant hole carriers, was evidenced by the positive S value observed across the entire temperature range.S value of the samples exhibited an initial ascent followed by a subsequent descent as the temperature rose, ultimately reaching its zenith at ~473 K. This behavior can be attributed to the influence of the bipolar effect [S of ~730 μVK−1 was obtained from the x = 0.6 solid solution. We calculated the Eg of the Cu3Sb(Sex1−Sx)4 (x = 0–1) samples with the formula: Eg = 2eSmaxT, where Eg, e, Smax, and T represent the band gap, elementary charge, maximal Seebeck coefficient, and the associated temperature, respectively [Eg for the pristine Cu3SbSe4 of ~0.30 eV aligned well with the reported literature [Eg from ~0.30 eV to ~0.69 eV, thus widening the band gap to alleviate the bipolar effect. For the semiconductors, we note that the increase in S (|S|) was directly proportional to the carrier effective mass and n−2/3. We calculated the Pisarenko relation between |S| and n  based on the single parabolic band model (SPB), as follows [Bk, ℏ represent the Boltzmann constant, and Planck constant, respectively. The calculated m* was significantly enhanced from 0.68 for pristine Cu3SbSe4 to 5.03 me for the x = 0.6 sample  of Cu3Sb(Sex1−Sx)4 (x = 0.1–1) samples were above the Pisarenko line. Furthermore, the m* depended directly on the Eg (E the energy of electron states), which deviated from a single Kane band model [S was related to Eg and m*. Consequently, the decreased carrier concentration (x > 0.5) and increased m*, resulted in the significant enhancement of S.Notably, the r effect . The maxectively . The calterature ,56; the  follows ,58:(1)S=6 sample . As seennd model ,59,60, tS2σ) of the Cu3Sb(Sex1−Sx)4 (x = 0–1) samples are presented in S2σ of Cu3SbSe4-Cu3SbS4 solid solutions exhibited a similar temperature-dependent behavior as the electrical conductivity (σ). The temperature-dependent trend observed in the S2σ was mirrored in the behavior of the σ for the Cu3SbSe4-Cu3SbS4 solid solutions. Owing to their relatively elevated σ and S values, these samples demonstrated larger S2σ values compared to the pristine Cu3SbSe4, particularly within the lower temperature range. Notably, the x = 0.3 sample achieved a larger S2σ value than the other samples, and the peak S2σ value of Cu3SbSe2.8S1.2 sample was ~670 μW m−1 K−2 at 673 K.The temperature dependence of the power factors (3Sb(Sex1−Sx)4 (x = 0–1) solid solutions are presented in κtot decreased with increasing temperature, mainly attributed to the increased scattering by lattice vibrations at elevated temperatures [κtot of pristine Cu3SbSe4 decreased from ~3.11 Wm−1K−1 at 300 K to ~1.03 Wm−1K−1 at 673 K. Similarly, the κtot of the x = 0.5 sample decreased from ~1.37 Wm−1K−1 at 300 K to ~0.52 Wm−1K−1 at 673 K. Generally, the κlat can be obtained by subtracting the electronic part (κele) from the κtot using the Wiedeman–Franz relationship (the details are displayed in L is the Lorenz number and it can be expressed as Equation (3) [The temperature dependence of the thermal transport properties of the Cueratures ,40,53 (Ftion (3) ,64:(3)L=L values of the Cu3Sb(Sex1−Sx)4 (x = 0–1) samples ranged from 1.5 to 1.6 W Ω K−2, and the results are listed in x < 0.6), the κele showed a slight increase at low temperature after S alloying, as described in κlat of the Cu3Sb(Sex1−Sx)4 (x = 0–1) samples is plotted in The calculated κlat, the κlat of the Cu3Sb(Sex1−Sx)4 (x = 0–1) compounds was evaluated at room temperature by the Debye–Callaway model. The primary scattering mechanisms under consideration were point defect scattering and Umklapp scattering. Then, the κlat of the pristine (κlat) Cu3SbSe4 compounds could be computed based on the Debye–Callaway model [u, θD, Ω, h, and ν represent the scaling parameter, Debye temperature, volume per atom, Planck constant, and average speed of sound, respectively  and Γs (strain field fluctuation) [x, ∆M/M and ∆r/r are the S concentration in one molecular, the relative change of atomic mass, and atomic radius owing to the replacement of Se with S, respectively. The ε value can be computed using the following formula [γ and υp are the Grüneisen parameter and Poisson ratio, respectively . Γ is ttuation) :(5)Γ=Γm+ formula :(6)ε=29( γ = 1.3  and υp =p = 0.35 ).m and Γs for the Cu3Sb(Sex1−Sx)4 compounds are presented in m and Γs changed with varying Se-alloying levels. It was observed that Γm was smaller than Γs when the S content x ≤ 0.6, indicating that the Γs (strain field fluctuation) contributed greatly to the drop of κlat. As for the x > 0.6 samples, the Γm (mass fluctuation) was the dominant. It is commonly accepted that the atomic radius of S is different from that of the Se atom, inducing a localized lattice distortion and leading to local field fluctuations that hinder the propagation of heat-carrying phonons [κlat closely aligned with the curve calculated by the Callaway model (κlat in the Cu3Sb(Sex1−Sx)4 solid solutions [κlat data with the recently reported values of Cu3SbSe4 are illustrated in 3Sb(Sex1−Sx)4 (x = 0.5) sample achieved an outstandingly low κlat of ~0.50 W m−1 K−1 at 673 K.The values of Γ phonons ,71. Howeay model e, suggesolutions . For a m3Sb(Se1−xSx)4 (x = 0–1) samples are illustrated in 3SbSe4. To further analyze our TE properties, the comparison of the ZTmax of the Cu3SbSe4-based materials is given in max of 0.72 was higher than that of other Cu3SbSe4-based materials, such as Cu3Sb0.97In0.03Se4 ~0.5, Cu3Sb0.985Ga0.015Se4 ~0.54, Cu3SbSe3.99Te0.01 ~0.62, Cu3Sb0.92Sn0.08S3.75Se0.25 ~0.67, Cu2.95Sb0.96Ge0.04Se4 ~0.70, and Cu2.95Sb0.98Sn0.02Se4 ~0.7 and is comparable to the ZT values for Cu3Sb0.98Bi0.02Se3.99Te0.01 ~0.76 and Cu3Sb0.91Sn0.03Hf0.06Se4 ~0.76. Although the Cu3Sb(Se1−xSx)4 (x = 0–1) samples had relative low ZT values in comparison with other high-performance TE materials, further enhancements of the ZT values can potentially be achieved by tuning the carrier concentration, dual-incorporation, and/or introducing band engineering.The temperature-dependent ZT of the Cu3SbSe4-Cu3SbS4 solid solutions were synthesized by vacuum melting and plasma-activated sintering (PAS) techniques, and the effects of S alloying on TE performance were investigated. S alloying can widen the band gap, effectively alleviating the bipolar effect. Additionally, the S (Seebeck coefficient) was significantly improved because of the increased m*. Furthermore, the substitution of S for Se in Cu3SbSe4 lattice led to noticeable local distortions, yielding large strain and mass fluctuations to suppress the κlat, thus decreasing the κlat and κtot to ~0.50 Wm−1K−1 and ~0.52 Wm−1K−1 at 673 K, respectively. Consequently, a peak ZT value of ~0.72 was obtained at 673 K for the Cu3Sb(Sex1−Sx)4 (x = 0.3) sample. Based on these results, it is speculated that further improvement in the figure of merit of Cu3Sb(Sex1−Sx)4 solid solutions can be obtained by enhanced electrical transport properties. Our research offers a new strategy to develop high-performance TE materials in solid solutions via elemental alloying.In summary, a series of Cu


How cells adapt their gene expression to nutritional changes remains poorly understood. Histone H3T11 is phosphorylated by pyruvate kinase to repress gene transcription. Here, we identify the protein phosphatase 1 (PP1), Glc7 as the enzyme that specifically dephosphorylates H3T11. We also characterize two novel Glc7-containing complexes and reveal their roles in regulating gene expression upon glucose starvation. Specifically, the Glc7–Sen1 complex dephosphorylates H3T11 to activate the transcription of autophagy-related genes. The Glc7–Rif1–Rap1 complex dephosphorylates H3T11 to derepress the transcription of telomere-proximal genes. Upon glucose starvation, Glc7 expression is up-regulated and more Glc7 translocates into the nucleus to dephosphorylate H3T11, leading to induction of autophagy and derepressed transcription of telomere-proximal genes. Furthermore, the functions of PP1/Glc7 and the two Glc7-containing complexes are conserved in mammals to regulate autophagy and telomere structure. Collectively, our results reveal a novel mechanism that regulate gene expression and chromatin structure in response to glucose availability. Chronic exposure to nutritional deficiencies or excess triggers metabolic stress, which is closely related to the incidence and progression of many chronic diseases1. Metabolic stress affects gene expression pattern, resulting in cellular and physiological changes3. To cope with environmental nutrient changes and maintain cell survival, cells need to reprogram their transcriptomes but the underlying mechanisms are not well understood5.Diet and nutrients have a profound impact on human development and health8. Cell metabolism can impact on chromatin modifications via the activity of metabolic enzymes and metabolites9. Several metabolites serve as substrates for chromatin modifications, such as acetyl-CoA for acetylation, S-adenosylmethionine (SAM) for methylation, and ATP for phosphorylation11. Other metabolites act as activators or inhibitors for chromatin modifying enzymes, such as α-ketoglutarate and NAD+11. Some metabolic enzymes even directly modify histones, which provides a new mechanism for specific regulation of gene expression by cell metabolism. The glycolytic enzyme, pyruvate kinase PKM2 has been reported to catalyze the phosphorylation of histone H3T11 (H3pT11) using phosphoenolpyruvate as its phosphate donor to promote gene transcription and tumorigenesis12. In budding yeast, pyruvate kinase Pyk1 translocates into the nucleus, where it phosphorylates H3T11 to repress gene expression within the SESAME complex (Serine-responsive SAM-containing Metabolic Enzyme)13. In addition, glycolysis and serine metabolism stimulate SESAME complex to phosphorylate H3T1114.Emerging evidence indicates that chromatin modifications are determined by cell metabolic status, which then affects gene expression and cell fates16. Autophagy is a conserved catabolic process that is initiated with the biogenesis of autophagosomes, which then fuse with the lysosomes or vacuoles to degrade their contents. Autophagy targets cytoplasmic contents, abnormal protein aggregates and damaged organelles for degradation and recycling to maintain cellular homeostasis and promote cell survival17. Under nutrient starvation conditions, autophagy is induced to a high level and cells need to activate the transcription of autophagy genes to replenish autophagy components19. The SESAME complex phosphorylates H3T11 at autophagy-related genes to repress their transcription and inhibit autophagy16. As SESAME-catalyzed H3T11 phosphorylation is enhanced with the increasing glucose concentrations, the autophagy activity is maintained at basal low levels under nutrient-rich conditions15.H3T11 phosphorylation has been reported to regulate the transcription of genes involved in autophagy15. Telomere has a typical heterochromatin structure and genes near these regions are subject to transcriptional repression, also known as telomere silencing21. Maintaining the integrity of telomere silencing is required for cell growth, development and differentiation24. Telomere heterochromatin is primarily maintained by assembly of the SIR (Silent Information Regulator) complex (Sir2/Sir3/Sir4), which binds at telomere regions and makes the underlying DNA inaccessible from the transcription machinery21. H3T11 phosphorylation enhances the SIR complex binding at telomere regions15. Moreover, H3T11 phosphorylation prevents autophagy-mediated degradation of Sir2, which further facilitates the binding of SIR complex at telomeres15. Although H3T11 phosphorylation directly connects autophagy with telomere silencing, how H3T11 phosphorylation is removed and dynamically regulated in response to nutritional changes remains largely unknown.H3T11 phosphorylation has also been reported to maintain gene silencing near telomere regions25. Glc7 has been reported to play important roles in regulating vesicle trafficking, cell polarity, DNA damage, transcription termination, cell-cycle progression, and etc.28. In this study, we identified histone H3T11 as a novel target for Glc7/PP1 and characterized two conserved Glc7-containing complexes that coordinately regulate autophagy and telomere silencing. Moreover, the activity of Glc7/PP1 to dephosphorylate H3T11 was enhanced under glucose starvation, resulting in induction of autophagy and reduction of telomere silencing. Most importantly, the above two Glc7-containing complexes are conserved in mammalian cells to regulate autophagy, telomere silencing and cellular senescence. Collectively, our study reveals that H3T11 phosphorylation is dynamically regulated by the highly conserved PP1 enzyme system, providing a mechanism by which cells regulate gene expression and chromatin structure in response to glucose availability.Glc7 is the sole catalytic subunit of protein phosphatase 1 (PP1), which was initially identified to regulate glycogen metabolism in budding yeast30. Among these protein phosphatase mutants, inactivation of PP1 phosphatase Glc7 in a temperature-sensitive mutant (glc7-12) at a non-permissive temperature (37 °C)31 significantly increased the intracellular H3pT11  that removes H3T11 phosphorylation (H3pT11), we first set out to screen the yeast protein phosphatase mutant library with anti-H3pT11 antibody. This library contains a total of 44 protein phosphatase mutants, including PPP family , PTP family , PPM family , RTR family , and HAD family T11 Fig. . No signT11 Fig. . The molT11 Fig. . To direT11 Fig. . SimilarT11 Fig. . As GLC715, we observed that Glc7 preferentially bound to gene coding regions, whereas H3pT11 was primarily enriched at upstream of the transcription start site (TSS) and downstream from the transcription end site (TES)  for Glc7. By comparing the ChIP-seq data for Glc7 and H3pT11ES) Fig. . Hence,  WT Fig. , furtherant Fig. , which m15. We first analyzed the ChIP-seq for the occupancy of Glc7 and H3pT11 at 39 autophagy-related genes. There was an anti-correlation between the occupancy of Glc7 and H3pT11 at ATG genes  was confirmed by ChIP-qPCR  for WT, H3T11A and ant Fig. . The dowPCR Fig. , indicatATG genes by dephosphorylating H3T11, we mutated H3T11A in glc7-12 mutant to construct glc7-12 H3T11A double mutant. The transcription of ATG genes was significantly reduced in glc7-12 mutant but was significantly increased in glc7-12 H3T11A mutant  and then switched to glucose starvation (SD-C) or nitrogen starvation (SD-N) medium. The intracellular H3pT11 levels were significantly reduced and the protein level of Glc7 was significantly increased when cells were grown in SD-C but not SD-N medium  medium and SD-C (glucose depletion) medium, respectively. The occupancy of Glc7 at gene coding regions was significantly increased when cells were shifted from SD to SD-C medium Fig. . The occium Fig. , which iium Fig. . In lineion Fig. ; howeverels Fig. . Nonetheium Fig. , which cced Fig. . Under gart Fig. . Mutatioant Fig. , suggestWhen cells were grown in SD-C medium, the occupancy of Glc7 at subtelomere regions was significantly increased as demonstrated by ChIP-seq Fig. . ChIP-qPglc7-12 mutant than WT upon glucose starvation  cells. Knockdown of PPP1CA significantly reduced LC3-II as well as the ratio of LC3-II/LC3-I in HeLa and primary human vein endothelial cells (HUVECs) cells  and knockdown of PPP1CA significantly increased the enrichment of H3pT11 at autophagy-related genes  and PPP1CA-knockdown  in mammalian cells. The Co-IP and reciprocal IP assay confirmed the interaction between PPP1CA and SETX Fig. . ChIP-qP49. Knockdown of PPP1CA significantly reduced the transcription of these three genes in HUVECs  staining, senescence-associated heterochromatin foci (SAHF) formation as visualized by DAPI stainingECs Fig. . MoreoveECs Fig. , indicatECs Fig. .Collectively, these data suggest that PPP1CA–Rif1 dephosphorylates H3T11 at subtelomere regions, which in turn impacts telomere silencing, telomere length and cellular senescence.15. However, it is poorly understood how H3T11 phosphorylation is removed and dynamically regulated in response to nutritional changes. In this work, we establish Glc7/PP1 as the enzyme that dephosphorylates H3T11 and identify two conserved Glc7-containing complexes as regulators of autophagy and telomere structure. When glucose is depleted, the expression of Glc7 is up-regulated, which then translocates into the nucleus to dephosphorylate H3T11 in the form of Glc7–Sen1 and Glc7–Rif1–Rap1 complexes. The Glc7–Sen1 complex dephosphorylates H3T11 to transcriptionally activate autophagy. The Glc7–Rif1–Rap1 complex dephosphorylates H3T11 to reduce telomere silencing. Interestingly, the Glc7–Sen1 complex accelerates autophagy-mediated degradation of Sir2, leading to reduced telomere silencing. Most importantly, both Glc7–Sen1 and Glc7–Rif1 complexes are conserved in mammals. PPP1CA interacts with SETX to dephosphorylate H3T11 and transcriptionally activate autophagy. PPP1CA interacts with Rif1 to dephosphorylate H3T11 at telomere-proximal regions, which then reduces telomere silencing, decreases telomere length and accelerates cellular senescence. Our results thus uncover a conserved mechanism by which PP1 dephosphorylates H3T11 to regulate gene expression and chromatin structure in response to nutrient availability.Cells need to coordinate gene expression with their metabolic states but the underlying mechanisms is not well understood. Pyruvate kinase phosphorylates histone H3T11 to repress autophagy and maintain telomere heterochromatin under nutrient-rich condition52. The eukaryotic organisms contain multiple genes encoding PP1 with the exception for budding yeast, which has a single gene, GLC7, that encodes PP153. This makes the yeast Glc7 as a good example to study the function and regulation of PP1. In this work, we identify histone H3T11 as the new target for Glc7/PP1 and reveal its role in transcriptional activation of autophagy. The genome-wide binding pattern for Glc7/PP1 anti-correlates that of H3T11 phosphorylation. As Pyk1 also binds at the coding regions in addition to gene promoter regions  and telomere-proximal genes. Our data showed that loss of Fbp1 had no significant effect on H3pT11 occupancy at these genes  can translocate into the nucleus to dephosphorylate histone H3T1127. The extensively studied Glc7 regulator is Reg1. Glc7–Reg1 dephosphorylates and deactivates Snf1 in the cytoplasm when glucose is abundant57. In addition, Glc7–Reg1 dephosphorylates and inactivates enzymes unnecessary for growth on glucose, such as Hxk258. However, loss of Reg1 has no effect on Glc7-mediated H3T11 dephosphorylation. Glc7 has been reported to be a subunit of the cleavage and polyadenylation factor (CPF) with Sen1, Nrd1 and Nab3 in the nucleus59. Loss of Sen1 significantly reduces the binding of Glc7 at autophagy-related genes, which increases H3T11 phosphorylation at autophagy-related genes. However, knockdown of Nrd1 and Nab3 has no significant effect on H3T11 phosphorylation, suggesting that Sen1 but not CPF complex is involved in recruitment of Glc7 to autophagy-related genes. Thus, Glc7–Sen1 could be a new complex that dephosphorylates H3T11 to activate autophagy. Although we cannot exclude the possibility that Glc7–Sen1 may directly regulate the transcription of autophagy-related genes, Glc7–Sen1-mediated H3T11 dephosphorylation could be a primary mechanism. As knockdown of Sen1 significantly increased the global H3pT11 level, the Glc7–Sen1 complex may not only dephosphorylate H3T11 at autophagy-related genes but also other genes. Indeed, by analyzing the ChIP-seq data for Sen1 and Glc7, we found that they co-localized at 73 genes, which are enriched in RNA degradation, fatty acid metabolism, autophagy, glycolysis pathways, and etc.Glc7 interacts with a multitude of regulators to control a diversity of processes such as glycogen metabolism, vesicle trafficking, cell polarity, DNA damage, gene transcription, and cell-cycle progression60. Our work provides a possible mechanism by which SETX activates autophagy by facilitating Glc7/PPP1CA-mediated H3T11 dephosphorylation. Interestingly, although Sen1 has no significant effect on telomere silencing under normal condition, Sen1 reduces telomere silencing by promoting Sir2 degradation under glucose starvation. This could be explained by the fact that under normal conditions, autophagy activity is low and H3T11 phosphorylation at subtelomere regions is relatively high, which restricts the nucleus export of Sir2 and protects Sir2 from being degraded by autophagy. Upon glucose starvation, autophagy activity is enhanced and H3T11 phosphorylation at subtelomere regions is reduced. Sir2 is dissociated from subtelomere regions, and translocates into the cytoplasm followed by degradation by autophagy. Under this condition, knockdown of Sen1 decreases autophagy activity, which may reduce autophagy-mediated Sir2 degradation and enhances telomere silencing.We identify a conserved role of Sen1 in regulating autophagy. Sen1/SETX recruits Glc7/PPP1CA to autophagy-related genes to dephosphorylate H3T11, which then activates the transcription of autophagy-related genes. Richard et al. reported that knockdown of SETX reduces autophagy flux but how SEXT regulates autophagy remains unknown61. Glc7 also acts as a Rif1 effector to suppress the double stranded break resection62. Sen1 has been reported to contain a “RVXF” Glc7 interaction motif (K1999SVAF2003) and mutation of F2003A abolished the interaction between Glc7 and Sen159. By analyzing the sequence of Rif1, we also identified a similar “RVXF” motif, K114SVAF118. It is possible that this “RVXF” motif is required for the interaction between Rif1 and Glc7. Moreover, the interaction between Rif1 with Glc7 is essential for suppression of TG repeat elongation28. Artificially tethered Glc7 can bypass the role of Rif1 in suppressing telomere extension28. However, it remains unknown about the target protein(s) of Glc7–Rif128. Here, we find that Rap1 recruits Rif1 to subtelomere regions, which then targets Glc7 to dephosphorylate histone H3T11 at subtelomere regions. Loss of Glc7–Rif1–Rap1 subcomplex increased H3T11 phosphorylation at subtelomere regions to enhance telomere silencing. Most importantly, the functions of Glc7 and Rif1 are conserved in mammalian cells: Rif1 directs PPP1CA to dephosphorylate H3T11 at subtelomere regions, which in turn reduces telomere silencing, shortens telomere length and accelerates cellular senescence. As deletion of RIF1 significantly increased the global H3T11 phosphorylation level, the Glc7–Rif1–Rap1 complex may not only dephosphorylate H3T11 at subtelomere regions. Like Glc7–Sen1 complex, the Glc7–Rif1–Rap1 complex may dephosphorylate H3T11 at other euchromatin regions enriched with Rif1.Rif1 has been reported to act as a Glc7 substrate-targeting subunit to prevent premature DNA replication origin firing by directing Glc7-catalyzed dephosphorylation of Minichromosome Maintenance complex28. Loss of Rif1 but not Rif2 increases telomere silencing63. Although both Rif1 and Rif2 can limit telomere elongation, Rif1 and Rif2 appear to act through different pathways as they have additive effects on telomere length41. One plausible explanation is that Rif1 and Rif2 may interact with different proteins once recruited to telomeres by Rap140. In our case, Rif1 but not Rif2 interacts with Glc7. It has been proposed that Rif1 regulate telomere length in part by competing with the Sir2/Sir3/Sir4 complex for binding to the C-terminus of Rap140. It remains unclear how Rif1 competes with the Sir2/Sir3/Sir4 complex. As H3T11 phosphorylation promotes the binding of Sir2/Sir3/Sir4 complex at subtelomere regions15, it is possible that Rif1 recruits Glc7 to dephosphorylate H3T11, which then dissociates the Sir2/Sir3/Sir4 complex at subtelomere regions. Our study hence provides a possible explanation about how Rif1 competes the Sir2/Sir3/Sir4 complex to bind Rap1.Although Rap1 recruits both Rif1 and Rif2 to telomere regions, only Rif1 affects Glc7-catalyzed H3T11 dephosphorylation. Loss of Rif2 has no significant effect on H3T11 phosphorylation. This differential behavior of Rif1 and Rif2 on telomere structure is consistent with the report by the Donaldson group, despite little is known about the underlying mechanism64. Oncogenic Ras promotes an increase in the intracellular levels of ceramides together with an increase in the PPP1CA protein level, which maintains pRb in a hypophosphorylation status64. Here, we report that knockdown of PPP1CA significantly reduced replicative senescence and uncovered a new mechanism by which PPP1CA promotes cellular senescence. PPP1CA dephosphorylates H3T11 at subtelomere regions, which reduces telomere silencing and shortens telomere length. Interestingly, a more recent study reported that PPP1CA is subjected to telomere silencing over long distance24. In senescent cells, PPP1CA expression is up-regulated, which could form a feedback loop to reduce telomere silencing and accelerate cellular senescence24. By directing PPP1CA to dephosphorylate H3T11 at subtelomere regions, Rif1 reduces telomere length and accelerates cellular senescence. Due to technical issues, it is difficult to obtain SETX-knockdown HUVECs and assess the effect of SETX on cellular senescence. Nonetheless, our study suggests that PPP1CA and Rif1 are potential targets for intervening premature cellular senescence.PPP1CA is a PPP1C isoform that contributes to the senescence program induced by oncogenic RasIn summary, we establish Glc7/PP1 as the enzyme that specifically dephosphorylates histone H3T11. Moreover, we identify two conserved PP1/substrate-targeting subcomplexes that regulate autophagy and telomere silencing. Our study uncovers a conserved mechanism to dynamically regulate H3T11 phosphorylation and reveals an epigenetic mechanism to regulate autophagy and telomere structure in response to nutrient availability.All yeast strains used in this study were listed in Supplementary Table /streptomycin solution. All cell lines used in this study were reauthenticated by short tandem repeat analysis after resuscitation in our laboratory.The primary HUVECs were purchased from ScienCell Research Laboratories. Cells were grown in Endothelial Cell Medium supplemented with 5% of FBS, 1% of Endothelial Cell Growth Supplement and 1% of penicillin/streptomycin solution. The HeLa and HEK293T cells were obtained from the American Type Culture Collection. The HeLa and HEK293T cells were maintained in Dulbecco’s modified Eagle’s medium supplemented with 10% FBS and 1% of penicillinYeast Extract Peptone Dextrose) medium at 30 °C until OD600 of 0.7–1.0. For glc7-12 mutant, cells were grown in 26 °C until OD600 of 0.5 and then transferred to pre-warmed medium at 37 °C for 2 h. For conditional knockdown mutants , cells were grown in YPD medium at 30 °C until OD600 of 0.5–0.7 and then treated with 40 μg/mL doxycycline for 2 h. Due to the lack of commercial antibodies against these proteins, the knockdown efficiency was examined by RT-qPCR. The knockdown efficiency results were presented in Supplementary Figs. 600 of 0.7 and then pelleted at 3000 rpm for 5 min. After being washed with sterile ddH2O, cells were grown in SD-N medium  for 0–6 h. For glucose starvation treatment, cells were grown in YPD until OD600 of 0.7 and then pelleted at 3000 rpm for 5 min. After being washed with sterile ddH2O, cells were grown in SD-C medium  for 0–6 h. For rapamycin treatment, cells were grown in YPD until OD600 of 0.7. 2 μg/mL rapamycin was added into the medium and cells were then grown for 0–1.5 h.Yeast cells were grown in 2% glucose-containing YPD . After centrifugation, the protein pellet was resuspended in 150 μL 2× SDS-sample buffer. Protein samples were separated by 8%–15% SDS-PAGE and transferred to PVDF membrane. The blots were probed by primary antibodies followed by incubation with horseradish peroxidase-labelled IgG secondary antibodies. The protein bands were visualized using the ECL Chemiluminescence Detection Kit  and quantified with Image J software (v.1.8.0).Cells were grown in 5 mL YPD or selective medium until ODAntibodies against H3 , H3pT11  and SETX  were purchased from Abcam; antibody against FLAG M2  was obtained from Sigma-Aldrich; antibodies against Sir2  and LC3  were purchased from Santa Cruz Biotechnology; antibodies against Myc , GAPDH , beta-actin , GFP  and goat polyclonal anti-mouse IgG  were obtained from Proteintech; antibody against Rif1  was purchased from Cell Signaling Technology; antibodies against CDKN1A/p21 , and PPP1CA  were purchased from Abclonal; antibody against CBP  was purchased from Absin. Antibody against Glc7 (1:5000) was custom-made in Abclonal and its specificity was confirmed by dot blots with Glc7 peptide . Chromatin was sonicated to an average size of ~500 bp and then subjected to IP with antibodies pre-bound to Protein G Dynabeads (Invitrogen) at 4 °C overnight. The beads were washed successively with FA buffer, FA buffer plus 1 M NaCl, FA buffer plus 0.5 M NaCl, TEL buffer  and TE buffer . The eluted DNA/protein complex was treated with 20 µg Proteinase K at 55 °C for 1 h followed by treatment at 65 °C overnight. After removing RNA by RNase (Roche), the DNA was purified with ethanol precipitation and quantitated by qPCR with primers listed in Supplementary Table For ChIP assay in yeast, cells were grown in 200 mL YPD media at 30 °C until OD65. Reads were aligned to yeast genome sacCer3 from UCSC using bowtie2 version 2.1.0 with parameter -k 1. The telomere repeated sequences were excluded. Data was put into R (3.1.0) for further analysis. The binding peaks were called using MACS2  with parameter -t -c -g 1.2e7 -n -B -q 0.01 --nomodel. Peaks annotation was performed on a website service (https://manticore.niehs.nih.gov/pavis2/). Tracks were smoothed by deepTools2 (v.2.0) and visualized by IGV software (v.2.0) with a reference genome (sacCer3).For ChIP-seq, the libraries were constructed and sequenced on an Illumina platform as described66. In brief, cells were cross-linked with 1% formaldehyde and quenched by 0.125 M glycine. After that, cells were collected, washed and lysed in buffer B . Chromatin was sheared by sonication and subjected to immunoprecipitation with antibodies pre-bound to Protein G Dynabeads (Invitrogen) overnight. Beads were washed and the eluted DNA/protein complex was then treated with 20 μg Proteinase K at 55 °C for 2 h and the crosslink was reversed at 65 °C overnight. After digestion with RNase (Roche), the DNA was purified and quantitated by qPCR with primers listed in Supplementary Table For mammalian cells, ChIP was performed as described15. Total RNA was extracted from HUVEC and HeLa cells by TRIzol reagent RNAiso Plus (Takara) as described66. Purified RNA was digested with DNase (Sigma) and reversed transcribed into cDNA using Reverse Transcriptase Kit (M-MLV) (ZOMANBIO). The RNA quality was determined by agarose gel electrophoresis. 0.5 µg RNA was taken for cDNA synthesis in NovoScript®Plus All-in-one 1st Strand cDNA Synthesis SuperMix (Novoprotein). The qPCR was performed using iTaq™ Universal SYBR® Green Supermix  with primers listed in Supplementary Table Total RNA was extracted from yeast cells by the phenol-chloroform extraction methodP value was calculated by edgeR (v3.24). The differentially expressed genes (DEGs) were defined as P < 0.05 and log2 fold change (FC) ≥ 0.75 or log2FC ≤ 0.75. Three biological replicates were performed for WT and glc7-12 mutant.Total RNA was isolated from exponential growing yeast cells by standard phenol-chloroform extraction procedures and the quality of RNA was examined using Agilent Bioanalyzer according to the manufacturer’s instructions. Library construction, sequencing and bioinformatics analysis were performed by MegaGenomics Inc. . The 65. The genes were grouped by their distance from the nearest telomeres in consecutive intervals of 10 kb. For each interval, the fraction of genes that were repressed in the mutant was calculated. A χ2 value for each interval was calculated by comparing the fraction of genes that were repressed in the interval to the genome-wide average. The statistical significance is defined as the following: χ2 > 3.841, P < 0.05; χ2 > 6.635, P < 0.01; χ2 > 10.828, P < 0.001.Genes within 60 kb of a telomere throughout the genome were pooled and ordered according to their distance from a telomere67. Yeast cells were cultured in YPD medium until OD600 of 1.0. After washing with cold PBS, cells were fixed with 4% formaldehyde in PBS for 0.5 h and treated with DAPI (Solarbio) for 15 min at room temperature. Cells were then washed with cold PBS. The cell morphology was visualized using a ZEISS LSM710 microscope (Germany) with a 100× oil immersion objective by fluorescent microscopy. Images were acquired using ZEN Imaging Software ZEN 2.1 (ZEISS). The merged color Images were generated by ImageJ.The microscopy and imaging processing were performed as described600 of 1.0, washed with PBS and flash frozen in liquid nitrogen. Thawed cell pellets were resuspended in IP buffer  and lysed with glass beads using a Biospec bead beater. The lysate was centrifuged at 45,000 rpm at 4 °C for 1.5 h and the supernatant was incubated with anti-FLAG agarose beads at 4 °C for 4 h. The beads were washed extensively in IP buffer. Proteins were disulphide-reduced by 25 mM DTT at 37 °C for 40 min, alkylated by adding 50 mM iodoacetamide, and then digested with sequencing-grade trypsin (Promega) at 37 °C overnight. The supernatant was desalted using C18 solid-phase cartridges and lyophilized. The dried peptides were reconstituted in 0.1% FA and loaded onto an Acclaim PepMap 100 C18 LC column (Thermo Fisher) utilizing a Thermo Easy nLC 1000 LC system (Thermo Fisher) connected to Q Exactive HF mass spectrometer (Thermo Fisher) and data were analyzed as described68.Yeast cells were grown in 1.2 L YPD media until an OD68. The crystal structure of Glc7 was obtained from the Protein Data Bank (PDB code 1ATP). To visualize the docked conformation, the PyMol molecular graphics system was used, which also removed water molecules and peptide inhibitor PKI. The structure of H3pT11 peptide (S10pT11G12) was generated using I-TASSER protein Structure and Function Prediction web server (http://zhanglab.ccmb.med.umich.edu/I-TASSER/). Molecular docking simulation of Glc7 and H3pT11 peptide was performed with the AutoDock Vina. All residues within Glc7 binding site were included using the following spatial coordinates of the central cavity: x = 126, y = 126, and z = 124. The coordinates of the grid resolution were x = 50.018, y = 50.565, z = 46.656.The structure modeling was performed as described2) at 30 °C for 0–1 h. The reaction was quenched by adding 2× SDS-sample buffer and boiling for 5 min. The reaction products were then subjected to immunoblots with anti-H3pT11 and anti-H3 antibodies.5 μg purified Glc7 was incubated with 100 ng recombinant purified histone H3 or in vitro-assembled nucleosomes in 30 μL dephosphorylation buffer 600 = 0.7–1.0) were collected, successively washed with ddH2O, PBS plus 10 mM β-mercaptoethanol and SB buffer  and then suspended in 1 mL SB buffer. 100 µL Zymolase 20 T (10 mg/mL in SB buffer) (mpbiochina) was added and samples were incubated at 30 °C with rotation until > 85% spheroblasts. Spheroblasts were collected by centrifugation at 2000× g, washed twice with SB buffer and suspended in 500 µL EBX . 0.5% Triton X-100 was then added to lyse the outer cell membrane. An aliquot was taken as the total fraction. The remainder of the lysate was layered over 1 mL NIB  and centrifuged at 13,000× g for 15 min. An aliquot of the upper layer cytoplasmic fraction was taken as the cytoplasm fraction. The glassy white nuclear pellet was suspended in 500 µl EBX and 1% Triton X-100 was added to lyse the nuclear membrane. Samples were kept on ice for 10 min and an aliquot taken as nuclear fraction.Yeast cells from 50 mL cultures , calmodulin Sepharose beads  or indicated antibody conjugated Protein A/G Sepharose beads  for 0–4 h at 4 °C. The beads were washed three times with IP washing buffer  and boiled with 2× SDS-sample buffer for immunoblot analysis.For Co-IP experiments in mammalian cells, cells were lysed in buffer A  at 4 °C for 0.5 h. After centrifugation, the supernatant was incubated with indicated antibody conjugated Protein A/G Sepharose beads  at 4 °C overnight, then washed with buffer A and boiled with 2× SDS-sample buffer for immunoblots.IFNB1) using qPCR as described70. The genome DNA (gDNA) was isolated using the animal genomic DNA kit  and quantified by spectrophotometer. 0.5–10 ng gDNA was used in qPCR with SYBR® Green Supermix . The primers used were:The relative human telomere length was quantified by measuring the amount of telomere repeat amplification product to a single copy gene product  was used for statistical analysis. *P < 0.05, **P < 0.01, ***P < 0.001, n.s., no significance. For all error bars, data represent mean ± SE.Representative results of at least two biological independent experiments were performed in all of the figure panels. The two-sided Student’s Supplemental Information


Such a catalyst is able to convert methane selectively to ethylene and propylene at a relatively low temperature (350 °C). The conception of the active center and construction of the catalyst have been described, and the origins of the catalytic performance are discussed.Nonoxidative coupling of methane exhibits promising prospect in that it affords value-added hydrocarbons and hydrogen with high atom economy. However, challenge remains in direct, selective conversion of methane to more valuable hydrocarbons like olefins. The current work presents a catalyst with well-dispersed Ta atoms anchored by graphitic C By contrast, direct methane conversion shortens the reaction paths and utilizes more proportion of methane + [3]+ [+ [4]+ +2]+ [3]+ , [TaN]+  [3]+ +  are able matters –25. The low 0.5% . The cur+ in the gas phase [4 center may possibly be built in the condensed phase. Here, we report a catalyst with Ta–N4 center as anchored in phthalocyanine that is supported by graphitic carbon nitride (g-C3N4). Such a catalyst is able to convert methane selectively to propylene at a relatively low temperature (350 °C).Notably, although Ta–O is an ideal center for methane activation, further conversion of the strong Ta–C bond may encounter high energy barrier –20. To was phase . Thus, ix center on the carbon-based materials. Thus, inspired by the synthesis method to prepare single atom catalyst [As the participation of oxygen in the transformation of methane is not wanted, we decided to construct an active Ta–Ncatalyst –33. Quit4 center was first reported in 1988 [4 center separately.The structure of tantalum phthalocyanine (TaPc) with a Ta–N in 1988 , and it  in 1988 ,35–38. T in 1988 . In fact in 1988 –42, the 3N4 has proven to be an ideal catalytic material for various chemical processes [3N4 was selected as the active center carrier in NOCM: (a) It is usually prepared above 500 °C through tubular furnace with high thermostability; (b) g-C3N4 is of 2-dimensional structure and high electron mobility, and it can support a structure like TaPc stably via π–π stacking [g-Crocesses . With thstacking –45; and stacking –47.3N4, the controlled sequential impregnation and pyrolysis procedures were found optimal. Through activation under hydrogen flow, the so-prepared TaPc/C3N4 catalysts exhibit surprisingly high activity toward methane conversion  values with over high Ta content. Thus, 0.08 wt% of Ta sample gives the highest TOF with the valve of 0.99 s−1 at 350 °C  sample. As shown in Fig. 3N4 heterojunction  sample; unfortunately, however, no useful information was obtained because of the rather low Ta content.To obtain more structural information of the active catalyst, the high-angle annular dark-field transmission electron microscope (HADDF-TEM) analysis was performed on the TaPc/Cion Fig.  reveals 3N4 material obtained via pyrolysis at 550 °C is thermally stable until 520 °C and completely decomposes at 750 °C. The TaPc complex undergoes a mass loss of 8.2% at 320 °C and is then kept stable until 550 °C; further decomposition occurs at 640 °C, leaving 11.52% ash eventually. The TaPc/C3N4 samples prepared at 400 and 500 °C have a good thermal stability below 500 °C, but the one prepared at 600 °C starts to lose mass at around 420 °C, probably because the backbone structure of the samples had already been destroyed during the preparation process. Considering the performances of different components of the TaPc/C3N4 sample under TGA, most likely, the impregnation–pyrolysis (<500°C) sequence retains the Ta–N4 structure of TaPc on the g-C3N4 surface.The thermal stabilities of different components of the catalysts were examined with thermogravimetric analyzer (TGA) analysis Fig. C. The g-3N4 diffraction peak at 27.5° offsets to 21.3° upon stacking the carbon nitride plane to form the TaPc/C3N4 (0.08 wt%) catalyst [3N4. Furthermore, there appear new diffraction peaks in the deactivated catalyst that are identified as C3N4  (00-053-0671) [0.83 (30.8° and 35.3°) (01-089-4765) [The x-rays diffraction patterns show that the typical g-Ccatalyst . This me53-0671)  and TaN089-4765) , respect3N4 (0.08 wt%) catalyst was probed with x-ray photoelectron spectroscopy analysis. As shown in Fig. 3N4 (0.08 wt%) spectrum has a red shift for Ta, which means that the electron density of Ta is increased, indicative for unstacked TaPc species.In addition, the electronic morphology of the TaPc/C4 units to the TaN clusters. Further, in the Raman spectrum of the deactivated catalyst, there is no peak at 1,300 and 1,600 cm−1 indicative for the D1 and G bands of aromatic carbon  supported on C3N4 (C90N123H15) via π–π stacking was thus built ; next, the so-formed ethyl group delivers a hydrogen to the hydride ligand of Ta to extrude molecular hydrogen (4 → 5), affording the N–CH(CH3)–Ta structure. The conversion of the third CH4 precisely resembles the 2 → → 5 sequence, thus proceeding along the 6 → → 9 path. Meanwhile, the formation of the N–CH(CH3)–Ta and N–C(CH3)2–Ta structures may compete with alternative channels. As shown in Fig. 2H4, C2H6, C3H6, and C3H8 molecules can be released from the reaction system via the sequences 2 → → 10/11 and 6 → → 12/13, respectively –Ta and N–C(CH3)2–Ta, all of them are entropically much less favorable as compared to the short paths 2 → → 5 and 6 → → 9. As a result, most likely, the N–CH(CH3)–Ta and N–C(CH3)2–Ta structures are the major intermediates the system goes through. However, these 2 structures may undergo further isomerization to release neutral olefin molecules: By shifting a hydrogen from the methyl group to the methylene carbon, the intact C2H4 or C3H6 is generated (2 → → 14 and 6 → → 15). Obviously, such 2 isomerization processes are highly competitive. As to the activation of a fourth CH4 molecule, although it is kinetically feasible at the structure 16, the lack of hydrogen in the bridging carbon prevents the conversion of the fourth methane from a sequential C–C coupling and H2 elimination process. Consequently, the propagation of the carbon chain stops at a number 3, and C3H6 is released as a major product. More details are provided in the Supplementary Materials. Although we tried to calculate as many as possible pathways, there might exist other various routes; after all, the generation of C2 and C3 involves multiple substrates as well as a series of bond breaking and making processes.According to the calculation, the most stable structure for the dehydrogenation of the first CH4 center anchored in phthalocyanine and supported on g-C3N4 was prepared. Such a catalyst is able to convert methane to ethylene and propylene at a relatively low temperature (350 °C) with the latter as the major product. A TOF value of 0.99 s−1 and >300-h one-pass lifetime prove the robustness of the catalyst. The deactivation results from the transformation of the Ta–N4 units to TaN clusters during the reaction process. Quantum chemical calculations indicate that the bridge N–CR2–Ta (R = H and CH3) structures serve as the key intermediates, which either enable the carbon-chain propagation or, alternatively, isomerize to release olefin molecules. Considering the vital role the Ta atom plays in the transformation that it serves as the courier to deliver both carbon and hydrogen, most likely, here, the relativistic effects induced strong Ta–C/H interaction functions once again. Future efforts may focus on how to uniformly load higher content of well-dispersed Ta on a carbon-based material and how to enhance the chemical stability of the Ta–N4 structure. Further, the relativistic effects exert similar influence on the gas and condensed phases, which encourages us to continue the gas-phase guided construction of high-performance catalyst. Indeed, gas-phase studies allow us to correlate various structural/electronic features with the performance of the active center, while the major difficulty is still on how these favorable features are replicated in bulk systems. As to the direct conversion of methane to higher hydrocarbons, most likely, a 5d element-centered structure is necessary so as to bind the carbonide intermediates strongly for further propagation of the carbon chain.In summary, a catalyst with single Ta–N3N4 was obtained.A mixture of cyanuric acid and melamine was heated to 550 °C with a rate of 5 °C/min in a tube furnace under Ar atmosphere, and the temperature was maintained at 550 °C for 5 h. After cooling to room temperature, faint yellow product of g-Cn-butanol was up to 100 °C, and a certain amount of DBU was next added. The system was stirred under reflux for 6 h in N2 atmosphere. The so-prepared reaction solution was mixed with g-C3N4 at N2 atmosphere and stirred for 24 h. After removing the solvent with vacuum evaporation, green solid was obtained. The dried green solid was calcinated in a tube furnace for 5 h to afford TaPc/C3N4. More technical details are provided in the Supplementary Materials.A mixture of phthalonitrile, tantalum pentachloride, and 2/90% He) in a quartz tube furnace. Next, the reactant flow (0.1% CH4/99.9% He) was fed in continuously under different temperatures. The product was analyzed using online gas chromatography. The reactivation of the catalysts was performed with calcination under reductive atmosphere (10% H2/90% He). For in situ IR analysis, the catalyst was placed in a reaction cell of the spectrometry and subjected to the same condition as that in methane conversion experiment. More technical details are provided in the Supplementary Materials.In catalytic performance, the catalyst was first activated under reductive atmosphere (10% H32N8H16Ta) supported on C3N4 (C90N123H15) via π–π stacking was built. Only the edge atoms of the C3N4 plane were frozen to avoid any boundary effects. More technical details are provided in the Supplementary Materials.The structural optimization and frequency analysis were performed at the GFN1-xTB level of theory using the xTB package (version 6.6.0) ,55 as in