Moreover, the formation of insoluble polysulfides  can slow down the reaction kinetics and terminate the discharge reaction before it reaches the final product. In this work, coffee residue derived activated ultramicroporous coffee carbon (ACC) material loading with small sulfur molecules (S2–4) as cathode material for RT Na–S batteries is reported. The first principle calculations indicate the space confinement of the slit ultramicropores can effectively suppress the formation of polysulfides . Combining with in situ UV/vis spectroscopy measurements, one‐step reaction RT Na–S batteries with Na2S as the only and final discharge product without polysulfides formation are demonstrated. As a result, the ultramicroporous carbon loaded with 40 wt% sulfur delivers a high reversible specific capacity of 1492 mAh g−1 at 0.1 C (1 C = 1675 mA g−1). When cycled at 1 C rate, the carbon–sulfur composite electrode exhibits almost no capacity fading after 2000 cycles with 100% coulombic efficiency, revealing excellent cycling stability and reversibility. The superb cycling stability and rate performance demonstrate ultramicropore confinement can be an effective strategy to develop high performance cathode for RT Na–S batteries.The formation of the soluble polysulfides (Na 2–4) as cathode material for room temperature (RT) Na–S batteries can effectively suppress the formation of polysulfides . Due to the special reaction mechanism, the ultramicroporous carbon loaded with 40 wt% sulfur (ACC‐40S) exhibits good rate capability, low‐self discharge, and excellent cycling stability.The coffee residue derived ultramicroporous carbon material loading with small sulfur molecules (S However, the scarcity of lithium resource leads to the increasing cost of lithium metal and makes Li–S batteries less attractive for grid‐related applications. In comparison, the abundant sodium resource on earth spur the strong interest to explore Na–S batteries as efficient energy storage technology for large scale grid application. Traditional Na–S batteries are usually operated beyond 300 °C with molten electrodes and solid β‐alumina (NaAl11O17) electrolyte, which induces severe safety concerns and additional costs, limiting their large scale applications. Therefore, room temperature (RT) Na–S batteries using liquid electrolyte attract great attention in recent years as low‐cost energy storage systems for large‐scale grid application. The major challenge for RT Na–S batteries is the nonconductive nature of S, which causes the low reactivity of S with Na and incomplete reaction to form polysulfides . The polysulfides dissolution and shuttling in the liquid electrolyte result in fast capacity fading, low coulombic efficiency, and large self‐discharge, restricting the practical application of RT Na–S batteries.To integrate the renewable energy sources such as wind and solar energies into large‐scale grids, it requires to develop highly efficient energy storage technologies with high energy density, long cycle life, and low cost.1 Amon As the solubility of sodium polysulfides is even higher than that of lithium polysulfides, finding an effective way to suppress sodium polysulfides formation is currently the top challenge for Na–S batteries. To improve the electrical conductivity and suppress polysulfides dissolution, various S host materials, such as carbon, conducting polymer, and metal oxides/sulfides, have been developed with composite cathode designs. Although significant progress has been achieved, the intrinsic problem of polysulfides formation cannot be completely ruled out and the electrochemical performance still cannot meet the requirements for large‐scale RT Na–S batteries. Among the carbon hosts, the design of microporous carbon was demonstrated to be very effective to improve the electrochemical activity of S and enhance the cyclability due to the micropore confinement. Two function mechanisms of micropore confinement have been proposed in previous works. By using molecular structure simulation with a cylinder pore model, Xin et al. believed that only Na2S2 and Na2S can be formed in the micropores due to the geometry confinement and small sulfur molecules (S2–4) in the microporous carbon. Wei et al. also found that only Na2S2 and Na2S can be formed in the microporous carbon, however, they attributed this behavior to a quasi‐solid‐state reaction because they believed that metal ions need to desolvate to enter the micropores. Nevertheless, previous works on microporous carbon indicate that the formation of soluble polysulfides can be effectively suppressed while the formation of insoluble polysulfides such as Na2S2 still exists for RT Na–S batteries. The insoluble polysulfides, as intermediate reaction products, have poor electronic conductivity and could slow down the reaction kinetics and impede the complete reaction to the final product of Na2S, thus reducing the reaction efficiency and specific capacity. Therefore, a one‐step reaction with Na2S as the only product and no polysulfide formation could be the ideal electrode design for RT Na–S batteries, which could simultaneously attain large specific capacity, good rate performance, and long cycle life. Such highly efficient electrode materials for RT Na–S batteries, however, have not been developed by far.As of now, sulfur loading of 70 wt% and above is viewed as the norm for Li–S batteries because Li–S batteries have been well studied in recent years and people now are working hard for practical applications. Na–S batteries, however, are still at the beginning stage of research now and the sulfur loadings in recent works on Na–S batteries are in the range of 40–60 wt%.9 As t2 g−1 was prepared from the coffee residue. By using density function theory (DFT) calculations, it is indicated that only small sulfur molecules (S2–4) can stay in the slit ultramicropores less than 0.6 nm. After infusing the sulfur into the micropores of the carbon, the carbon–sulfur composite electrode with 40 wt% sulfur loading (ACC‐40S) exhibited only one voltage plateau for both charge and discharge and further in situ UV/vis spectroscopy measurements demonstrated Na2S is the only product during the discharge. The DFT calculations indicate that Na2S2 molecules cannot keep its corrugated configuration in the micropores less than 0.7 nm due to the space confinement effect, while the Na2S molecules with the planar structure can be formed without distortion in micropores less than 0.5 nm. By using the ACC‐40S as cathode, the RT Na–S batteries exhibited a high specific reversible capacity of 1492 mAh g−1 at 0.1 C and excellent cycling stability with nearly no capacity fading for 2000 cycles at 1 C. Importantly, the RT Na–S batteries using the ACC‐40S cathode also exhibited low self‐discharge with a low capacity‐fading rate of 0.17% per day, making them promising for potential applications in large‐scale energy storage.Herein, we report a micropore confinement cathode that enables one‐step reaction without formation of polysulfides for RT Na–S batteries. The microporous carbon with a major micropore size around 0.5 nm and high surface area of 1100 m28 molecules to break into the metastable S2–4 small molecules and migrate into the ultramicropores. The space‐confinement effect of ultramicropores not only limits the volatilization of S2–4 at high temperatures, but also inhibits their recovery to S8. To understand why only small sulfur molecules can exist in the ultramicropores, the adsorption energies of S2–5 and S8 molecules inside ultramicropores were calculated by using a bilayer graphene model to simulate the slit ultramicropore. The optimized structures of S2–5 and S8 molecules inside the bilayer graphene with different interlayer spacings are shown in Figure 2–4 could stably exist inside the bilayer graphene, and the buckled S5 and S8 molecule structures will finally break into S2 and S3. The adsorption energies of the various S‐molecules inside the bilayer graphene are summarized in Table S1 (Supporting Information). The negative energy means the insertion of an S molecule is an exothermal reaction, which is thermodynamically feasible. It is found that the insertion of S5–8 molecule into the bilayer graphene can reach a stable state only when the interlayer spacing is larger than 0.7 nm. In contrast, the more positive of the adsorption energy indicates the smaller possibility of the existence of S‐molecules. Therefore, if the micropore size of the microporous carbon can be controlled less than 0.7 nm, only small S2–4 molecules can stably exist in the slit micropores.The facile preparation of ACC and the process of immobilizing sulfur into the carbon ultramicropores are schematically illustrated in cropores.12 TheBy controlling the sulfur infiltration, the carbon–sulfur composites with various sulfur loadings of 40, 50, and 60 wt% were successfully obtained and denoted as ACC‐40S, ACC‐50S, and ACC‐60S, respectively. Figure S2a,b (Supporting Information) show the photo images of the ACC‐40S sample before and after heating. It is obvious that the color changes from gray to black after heating, which suggests the infiltration of sulfur into carbon. Figure S3a (Supporting Information) shows the field emission scanning electron microscopy (FESEM) image of the pristine carbonized coffee residue (CC) without any further treatment, revealing a relatively dense structure without obvious pores being observed. With KOH activation, the obtained ACC sample as shown in Figure S3b (Supporting Information) presents a 3D porous structure with microsized pores interconnnected. In this 3D porous structure, the surface area of the ACC can be greatly increased, and most of the micropores in the carbon can be easily accessed by the small sulfur molecules. Figure S3c (Supporting Information) shows the FESEM image of the final obtained ACC‐40S sample. The well preserved 3D structure and the smooth surface indicate all of the sulfur is infused into the micropores of the carbon without extra sulfur on the carbon surface. The Figure S4a,b (Supporting Information) show the FESEM images of the ACC‐50S and ACC‐60S samples, respectively, which reveal rough carbon surface, suggesting extra sulfur on the carbon surface when the sulfur content is exceeding 40 wt%. Except for the weak diffraction peaks from carbon, the XRD pattern of ACC‐40S does not show any diffraction peaks of S8, demonstrating only small sulfur molecules are confined in the micropores without large S8 molecules at surface. In contrast, the diffraction peaks of S8 phase (JCPDS No. 85‐0799) can be detected in the XRD patterns of the ACC‐50S and ACC‐60S, indicating that only about 40 wt% sulfur can be loaded in the micropores and further increased sulfur content will result in large S8 molecules at the surface without micropore confinement. In addition, the S8 diffraction peaks still cannot be detected in the ACC‐40S electrode after several charge/discharge cycles, as shown in Figure S5 (Supporting Information), demonstrating the sulfur confinement in the micropores is very stable. The Raman spectra of CC, ACC, and ACC‐40S were further compared in Figure ID/IG) is an indicator of the defect degree in carbonaceous materials. As shown in Figure ID/IG ratio of ACC (0.94) is much higher than that (0.86) of CC due to the structure distortion induced by the activation process, which is consistent with the XRD results. It is worth noting that the two peaks located at 176 and 805 cm−1 are observed in the Raman spectrum of ACC‐40S, which could be ascribed to the chemical bonds between sulfur and carbon. To further investigate the chemical constitutions of CC, ACC, and ACC‐40S samples, Fourier transform infrared spectroscopy (FTIR) measurements were carried out and the obtained FTIR spectra are shown in Figure −1 can be ascribed to C–S group, which is in good agreement with the Raman results.X‐ray diffraction (XRD) patterns of CC, ACC, ACC‐40S, ACC‐50S, and ACC‐60S are shown in 13 Exc respectively, suggesting the N‐containing groups on the carbon surface and the chemical bonding between sulfur and carbon. Such interaction between sulfur and carbon could be beneficial to effectively immobilize the sulfur in the carbon micropores. Figure 3/2 and 2p1/2, respectively. The slightly lower binding energy of S 2p3/2 of ACC‐40S than that of pure S powder suggests the interaction between sulfur and carbon, which is consistent with Raman and FTIR results. Furthermore, a minor peak at 168.9 eV is ascribed to the sulfate species formed by partial oxidation of sulfur in air. The N 1s core‐level XPS spectrum is shown in Figure S7a (Supporting Information), in which three components can be assigned to pyridinic N (399.7 eV), pyrrolic N (400.9 eV), and graphitic N (402.8 eV), respectively. The high‐resolution XPS spectrum of O 1s in Figure S7b (Supporting Information) exhibits three components at 531.8, 532.6, and 533.8 eV, which can be assigned to C═O, C─O, and ─OH species, respectively. After the ACC‐40S electrode is fully discharged, a new peak located at 290 eV emerges in the core‐level XPS spectra of C 1s , the core‐level XPS spectra of N 1s shows similar peaks comparing to the ACC‐40S but a new Na Auger peak appears at the discharge state in the high‐resolution XPS spectrum of O 1s. For comparison, the core‐level XPS spectra of O 1s, N 1s, and C 1s of CC and ACC are shown in Figures S7 and S8 (Supporting Information), which are similar to those of ACC‐40S. To investigate the interaction between Na2S and each functional group, we add one type of functional group in monolayer graphene model every time separately, as shown in Figure S9 (Supporting Information). Here, we added the optimized models of functionalized graphene, and calculated the adsorption energies of Na2S on pristine and functionalized graphene . It is found that the functional groups on graphene have strong interaction with Na2S molecule by forming O─S, N─S bonds. Especially, if the ─OH can stay inside of the pores and Na2S molecule is put nearby, the ─OH will detach from the carbon surface and form a new product of Na2SOH . However, such O─S, N─S bonds, and Na2SOH (S−1) cannot be detected in our XPS results, indicating such functional groups are negligible inside the pores. On the other hand, the existence of C─O and C─OH bonds can induce the lattice distortion of carbon and cause the surface roughness from 0.6 to 1.1 nm, which is not feasible inside the micropores with major pore size of 0.5 nm for the present work. To rule out the effect of coffee origin on the experimental results, we collected some coffee residue from USA and prepared the porous carbon by using the same synthesis procedure. As shown in Figure S10 (Supporting Information), the obtained carbon sample from USA coffee residue exhibits the same microporous structure from BET results and similar N, C, O compositions from XPS results. The thermogravimetric analysis (TGA) of ACC‐40S in Figure X‐ray photoelectron spectroscopy (XPS) measurements were carried out to investigate the chemical species of the CC, ACC, and ACC‐40S samples. As shown in Figure S6a,b (Supporting Information), the XPS survey scan spectra indicate the presence of C, N, and O elements in both CC and ACC samples. Four elements of C, N, O, and S can be detected in the XPS survey scan spectrum of ACC‐40S and Na signal can be further detected when the ACC‐40S electrode is fully discharged . As shown in Figure 18 ress Figure , which cs Figure  shows an2 g−1 by using the Brunauer−Emmett−Teller (BET) method. It is noted that the adsorption increases sharply when p/p0 is below 0.02 (due to micropore filling) and then the adsorption curve becomes nearly horizontal, which indicates the ultramicropores are slits as illustrated in Figure  A hysteresis loop at p/p0 above 0.95 associates with the presence of large mesopores arising from textural mesoporosity. The micropore size distribution and the cumulative pore volume distribution calculated by Horvaih–Kawazoe (HK) method is presented in Figure 3 g−1 for ACC, which corresponds to a theoretical loading of 44 wt% sulfur in micropores. Except for micropores, widely distributed secondary mesopores ranging from 30 to 200 nm could be observed in the pore size distribution curve . As shown in Figure 2–4 are well dispersed in the ultramicropores of ACC.The pore structure of ACC is further characterized by nitrogen adsorption/desorption isotherms. −1 between 0.5 and 2.9 V for the initial three cycles are shown in + diffusion into the micropores. Different from previous works, there is only one cathodic peak for all cycles, corresponding to the direct formation of Na2S. For the anodic scan, the peak at around 1.9 V can be assigned to the reverse oxidation process to extract Na+ from the electrode. The CV curves of the ACC‐40S electrode at different scan rates of 0.1, 0.3, 0.5, and 1 mV s−1 are revealed in Figure S13 (Supporting Information). The reduction peaks are seen to shift toward lower voltage while the oxidation peaks move toward higher voltage with increasing scan rate, which can be to the increased polarization for the cell at high scan rates.The electrochemical properties of RT Na–S batteries using ACC‐40S as cathode were evaluated through the CR2032‐type coin cells. The cyclic voltammetry (CV) curves of the ACC‐40S electrode at a scan rate of 0.1 mV scropores.30 Dif−1 for the first cycle and can still remain 1110 mAh g−1 after 400 cycles. Even when the electrode loading of the ACC‐40S electrode was increased up to 3 mg cm−2, it can still deliver a high reversible capacity of 960 mAh g−1 at 0.1 C after 200 cycles . The discharge/charge curves and cycle performances of the ACC‐xS  electrodes with high sulfur loadings are shown in Figure S15 (Supporting Information) for comparison. As the micropores in the current carbon material can only accommodate about 44 wt% sulfur, extra sulfur loading can proceed through the carbon surface adsorption. However, the surface adsorbed sulfur is easy to get lost due to polysulfide formation and dissolution during the charge/discharge processes. Accordingly, the ACC‐50S and ACC‐60S electrodes show much lower specific capacities of 816 and 705 mAh g−1, respectively, after 400 cycles. However, with a high sulfur loading of ≈70 wt%, the ACC‐70S electrode shows a visible plateau at around 1.9 V during the first discharge process, indicating the generation of polysulfides. When the sulfur loading reaches 80 wt%, such high voltage plateau is more pronounced in the first discharge curve, suggesting severe formation of polysulfides. Obviously, the cell performance is significantly deteriorated by the high sulfur loading and formation of massive polysulfide. Both ACC‐70S and ACC‐80S electrodes exhibit very low specific, indicating very low utilization of sulfur in the electrodes. For further comparison, the CC‐40S and commercial activated carbon–sulfur (CAC‐40S) electrode were also tested under the same conditions . The charge/discharge curves of both the CC‐40S and CAC‐40S electrodes show very low reversible capacities, indicating very low utilization of sulfur in both electrodes.Figure rocesses.32 Acc The photo images of the separators in the cells using CAC‐40S (the first row) and ACC‐40S (the second row) as cathodes after 0, 10, and 100 cycles are shown in Figure Since the separator is in close contact with the cathode, the color change of the separator during cycling can be used as an indicator, reflecting the formation and dissolution of sodium polysulfides from the cathode.34 The−1 at 0.1, 0.2, 0.5, 1, 2, and 5 C respectively, demonstrating good rate capability. Notably, when the current density returns to 0.1 C, a reversible capacity of 1170 mAh g−1 could still be recovered, indicating high reversibility for the electrode. The superior rate performance of the ACC‐40S electrode can be attributed to its unique 3D structure, facilitating electrolyte penetration, and small sulfur molecules confined in the micropores, favoring fast reaction kinetics. It is worth noting that the ACC‐40S electrodes exhibits superb cycling stability even at high current densities. As shown in Figure To further investigate the rate performance, the charge/discharge measurements of the ACC‐40S electrode were carried out at different current densities from 0.1 to 5 C. As shown in Figure  As reference, the UV/vis spectra first‐order derivative curves of different Na2Sx (2 ≤ x ≤ 8) products are displayed in Figure 2S can be detected without any peak shifting during the whole discharge process, demonstrating that Na2S is the only product during discharge. The in situ UV/vis spectroscopy can detect the dissolved Na2S and polysulfides in the electrolyte. The electrolyte we used in this study is 1.0 m NaClO4 in ethylene carbonate (EC)/propylene carbonate (PC), and the solubility of Na2S in this electrolyte is very low as compared to those of the dissoluble polysulfides. Due to the very low solubility of Na2S, the UV signal of Na2S dissolved in the electrolyte does not change much, and cannot reflect the variation of Na2S in the electrode. Neither soluble polysulfides nor insoluble polysulfides can be formed during the discharge process in the ACC‐40S electrode, which agrees well with electrochemical behavior, indicating one‐step redox reaction for the present Na–S system. In order to ensure the reversibility of the cell, in situ UV/vis spectra for the cell were also measured at different voltages during the charge process , which indicates sulfur can be the only charge product without polysulfides formation. Therefore, it can be confirmed that no polysulfides  can be formed during the whole reaction process for ACC‐40S electrode, and this new one‐step reaction mechanism could be the main reason resulting in the ultrastable cycling for the RT Na–S batteries.In order to investigate the reaction mechanism and the reaction products during battery cycling, in situ UV/vis spectroscopy measurements were performed at different discharge states as shown in 35 As 2S2 molecule has a 3D structure, while the Na2S molecule is 2D, and this difference is the key factor that makes Na2S easier to form in the slit micropores. The possible reaction mechanism for converting small sulfur molecules Sn (n = 2–4) into Na2S is proposed and shown in Figure S20 (Supporting Information). Figure 2S and Na2S2 molecules intercalated into the bilayer graphene (simulating the slit micropore) with different interlayer spacings. Na2S with a plane structure can be laying in the narrow slit even when the interlayer spacing is less than 0.5 nm without changing its molecular configuration. Since the interaction between Na2S and the carbon‐layer (up and down) is van der Waals (vdW) force, from Figure 2S to the up/down carbon layer is about 0.3 nm, which is a typical distance for vdW interactions. With the decreasing of interlayer spacing, the repulsion between Na2S and carbon layers increases, which leads to the enhancement of total energy. When the interlayer spacing is larger than 0.60 nm, the total energy of the system also slightly increases, because the increasing of distance between Na2S and up carbon layer weakens the interaction, which indicates Na2S molecules are free to form in the interlayer when interlayer distance is larger than 0.60 nm.To further understand the one‐step reaction of the present Na–S battery without polysulfides formation, the reaction mechanism of the ACC‐40S electrode with geometry consideration for micropores is illustrated in Figure 2S2, a larger interlayer distance is required to keep its corrugated configuration. Seen in Figure 2S2 molecule will be distorted when the interlayer spacing is less than 0.75 nm. As shown in Figure 2S2 could form when the interlayer distance is larger than 0.75 nm. Considering the micropore size distribution is dominated at about 0.46 nm and the pore size distribution above 0.75 nm is nearly zero in our ACC, the Na2S2 molecules are hardly to form in the slit micropores from the thermodynamic analysis, agreeing well with our experimental results. Therefore, as a result of micropore confinement, Na2S is the only product for the Na–S system, enabling the one‐step reaction mechanism without polysulfides formation. The optimized structures of Na2S and Na2S2 molecules intercalated into the bilayer graphene with other interlayer spacings are also shown in Figure S21 (Supporting Information). Although Xin et al. thought that both Na2S2 and Na2S molecules can be accommodated as single molecule in the cylindrical micropores with 0.5 nm diameter, the difference may be induced by the effect of curvature, because they used a cylindrical micropore model.In case of Na2S molecules still can exist, but the charge transfer from Na2S to carbon layer has been enhanced and the probability to form Na─C/S─C bonds are raised. The partial the density of states (PDOS) of the systems with one Na2S molecule in narrow slit with 0.5 and 0.8 nm distance are shown in Figure 2S maintains a molecular structure. The interaction between Na2S and slits is van der Waals forces. While for Figure 2S during discharging. The strong interaction between Na2S and carbon layer improves the charge transfer from Na2S to carbon layer and enhances the carrier concentration of carbon layer. As the electronic conductivity is proportional to carrier concentration, the confinement of Na2S molecules into the slit micropores is beneficial to improve the electron transfer, thus further increasing the reaction efficiency.When the interlayer spacing is less than 0.45 nm, which is close to bonding distance, the strong repulsion causes sharply increased total energy of the system. When the interlayer distance ranges from 0.45 to 0.60 nm, Na32–4) confined in the slit micropores by tuning the infused sulfur content. The electrochemical behavior of the carbon–sulfur composite electrode with 40% sulfur content (ACC‐40S) indicates one‐step Na–S reaction, and the in situ UV/vis spectroscopy result demonstrates Na2S is the only discharge product without formation of polysulfides during charge/discharge for the Na–S batteries. As a result, the ACC‐40S electrode exhibits a high reversible specific capacity up to 1492 mAh g−1 at 0.1 C, good rate capability, low‐self discharge, and excellent cycling stability with negligible capacity loss at 1 C after 2000 charge/discharge cycles. The micropore confinement has been elucidated with the DFT calculation results using slit micropore model, demonstrating formation of solo Na2S product in the composite electrode is thermodynamically feasible by tailoring proper micropore size. The present work provides new insights for micropore confinement and porous carbon materials design for developing ultrastable RT Na–S batteries without polysulfide formation.In summary, we report the preparation of coffee residue derived 3D microporous carbon with small sulfur molecules  in a mortar for at least 20 min. Then the mixtures were transferred into Teflon containers and preserved at 155 °C for 15 h. At last the black powder was heated at 200 °C for 20 min in Ar atmosphere to obtain the ACC‐xS. For comparison, the CC‐40S and CAC‐40S were also prepared through the same preparation process.The different ACC‐The morphologies and microstructures of the samples were investigated by FESEM (Quant 250 FEG) and HRTEM (Tecnai G2 F30 S‐TWIN). Nitrogen adsorption and desorption isotherms were measured with an adsorption apparatus (JW BK‐200C). The crystal structure and phase purity of the as‐prepared samples were characterized using XRD (Bruker‐AXS D8 Advance) and Raman (Jobin‐Yvon T6400 Micro‐Raman system with laser excitation length of 514.5 nm). The sulfur content of the composite was measured by TGA (SDTA851E). The chemical composition and electronic structure of the samples were investigated by XPS (ESCALAB250Xi with a monochromatic Al anode X‐ray source) and FTIR with a Nicolet FTIR Is 10 spectrometer.N‐methyl pyrrolidinone with a mass ratio of 8:1:1 to form a homogeneous slurry, and then the slurry was coated on an aluminum foil. The electrodes were dried at 70 °C for 24 h in a vacuum oven. The loading density of the sulfur was about 1 mg cm−2. To test the electrochemical performance, the cells were assembled in an argon‐filled glove box in which the water and oxygen concentrations are kept below 0.1 ppm. Here the ACC‐xS, CC‐40S, and CAC‐40S were used as working electrodes, Na foil as counter and reference electrode, glass fiber as separator (Whatman GF/F), and 1.0 m NaClO4 in EC/(PC) (1:1 by volume) with 2 wt% fluorinated ethylene carbonate (FEC) as electrolyte. The CV measurements were carried out using a CHI660D electrochemical workstation. Charge/discharge measurements were performed in a potential range of 0.5 and 2.9 V (vs Na+/Na) at different current densities using LAND CT2001A battery tester.To prepare the electrodes, active materials, super‐P, and poly acrylic acid (PAA) binder were mixed in 2Sx . The sodium sulfide and sulfur with molar ratios of 1:0, 1:1, 1:2, and 1:3 were dissolved in deionized water via magnetically stirring for 2 h to synthesize Na2S, Na2S2, Na2S3, and Na2S4, respectively. And the Na2S6 and Na2S8 were synthesized with the sodium sulfide to sulfur molar ratios of 1:5 and 1:7 dissolved in TEGDME. The adsorption ability of Na2Sx were qualitatively determined using the UV–vis spectrometer and the effect of the solvent would be removed during the testing process.Six different sodium polysulfide solutions with varied ratios of sodium to sulfur were prepared, which are denoted according to the sulfur content in Nam NaClO4 in EC/PC (1:1 by volume) with 2 wt% FEC was used to infiltrate the separator. As the cathode material, the sulfur loading of the ACC‐40S was about 1 mg cm−2. At last, the cell was assembled in a traditional way.The in situ UV–vis spectroscopy was operated on UV Lambda 750S UV/Vis/NIR spectrometer. The positive case and the sodium metal were punched and the holes were half size of the cathode. 1.0  based on DFT. The generalized gradient approximation (GGA) with Perdew–Burke–Ernzerhof (PBE) algorithm is selected as exchange‐correlation interaction. To improve the description of long‐range van der Waals interactions, the DFT‐D3 method was employed. An energy cutoff of 500 eV and a Monkhorst–Pack scheme with a k‐point mesh of 11 × 11 × 1 were employed to ensure the calculation accuracy. The structures of bilayer graphene were fixed and Na2S and Na2S2 molecules were relaxed with energy change converged to 1 × 10−5 eV in electronic self‐consistency cycles and all force smaller than 0.02 eV Å−1 in ionic relation loops. The PDOS were obtained by the Atomistix ToolKit (ATK) package. The geometry optimizations are performed by the Vienna Ab initio Simulation Package (VASP), 38 basThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


Flexible metal–organic frameworks (MOFs) show large structural flexibility as a function of temperature or (gas)pressure variation, a fascinating property of high technological and scientific relevance. The targeted design of flexible MOFs demands control over the macroscopic thermodynamics as determined by microscopic chemical interactions and remains an open challenge. Herein we apply high‐pressure powder X‐ray diffraction and molecular dynamics simulations to gain insight into the microscopic chemical factors that determine the high‐pressure macroscopic thermodynamics of two flexible pillared‐layer MOFs. For the first time we identify configurational entropy that originates from side‐chain modifications of the linker as the key factor determining the thermodynamics in a flexible MOF. The study shows that configurational entropy is an important yet largely overlooked parameter, providing an intriguing perspective of how to chemically access the underlying free energy landscape in MOFs. By combining high‐pressure powder X‐ray diffraction and molecular dynamics simulation we gain insight into the microscopic chemical factors that determine the high‐pressure macroscopic thermodynamics of two flexible pillared‐layer MOFs, identifying configurational entropy as originating from side‐chain modifications as the determining factor of the macroscopic thermodynamics. At the centre of applied inorganic chemistry and materials science is the development of structure‐property relationships and the search for materials with (tailored) physicochemical properties of scientific and technological relevance. In this context metal‐organic frameworks (MOFs) have proved as a tantalizing material platform, providing a nearly unlimited parameter space for exploring a wide range of properties such as tuneable nonlinear optical properties,V=20 % as response to temperature and pressure variation, and guest adsorption.[2, with im−=imidazolate)[2+=Fe2+ or Co2+, bdp2− = 1,4‐benzenedipyrazolate)[A subclass of MOFs, so‐called flexible MOFs, shows large structural flexibility with volume changes exceeding Δsorption.5, 6 I3+=Al3+, Fe3+, Cr3+, bdc2−=1,4‐benzenedicarboxylate),[lp) to narrow pore (np) phase transition. Importantly, tunability of the transition pressure was recently shown by Yot et al., who implemented functional groups (‐Br and ‐CF3) into the MIL‐47(V) framework,p<0.1 GPa),[More generally, flexible MOFs offer the opportunity for the creation of smart materials with distinct responsiveness towards different external stimuli. In this context hydrostatic pressure as external stimulus is gaining more and more attention,23, 242(fu‐bdc)2dabco .[2− linkers form 2D sheets with sql topology, see Figure pcu topology. It has been shown that the temperature and gas sorption behaviour of fu‐MOFs can be controlled via functionalisation of the bdc2− linker which can render the MOFs flexible.[2(BME‐bdc)2dabco (BME=bis(methoxyethoxy)) contracts to a np phase after guest removal and shows a np to lp phase transition at T=493 K.cf. Figure 2(fu‐bdc)2dabco series and in flexible MOFs in general.[2(fu‐bdc)2dabco to hydrostatic pressure is yet entirely unexplored, providing us with another fascinating angle from which to probe the free energy landscape of these materials as a function of chemical changes.Whilst recognizing the increasing number of high‐pressure studies that have appeared in recent years, it can be observed that the overall number of such studies is still limited, arguably due to experimental limitations.39, 402(DB‐bdc)2dabco  and compare it to the properties of the parent material Cu2(bdc)2dabco by combining state‐of‐the‐art high‐pressure powder X‐ray diffraction (HPPXRD) and atomistic resolved force field molecular dynamics simulations. Often fu‐MOFs undergo a (reversible) lp to np phase transition upon guest removal from the pores,[2(DB‐bdc)2dabco remains in the lp phase,K), the transition pressures and the onset of amorphisation processes. The experiments are complemented with molecular dynamics simulations, monitoring the phase transitions and assessing the underlying energetics computationally. The simulations are based on our recently re‐parameterized MOF‐FF force field (MOF‐FF) for the accompanied side chain library,2(DB‐bdc)2dabco as the first example in which configurational entropy dictates the observed physicochemical properties.In this work we explore the high‐pressure response and thermodynamics of Cuhe pores,41, 422(bdc)2dabco and Cu2(DB‐bdc)2dabco. From this data, the presence of phase transitions and amorphisation processes is observed by visual inspection. In a subsequent quantitative analysis, the evolution of the unit cell volume as a function of pressure is extracted from which the bulk moduli are obtained. These experimental data are used for benchmarking the outcomes of the molecular dynamics simulations, guiding us in the analysis and interpretation of the results. The computational approach is based on the protocol introduced by Rogge et al.,p(V) equation of state and the thermodynamic properties along the volume change from the lp to the np form. Since we recently observed a strong dependency of the outcome of such simulations on the initial configuration of the alkoxy side chains in materials such as Cu2(DB‐bdc)2dabco and Zn2(DB‐bdc)2dabco,lp to np phase transition which is only accessible through molecular dynamics simulations due to the large number of structural distortions such as various number of side chain conformations.We start by analysing the experimental HPPXRD data of Cup=ambient−0.4 GPa (Δp per step=0.025–0.05 GPa) on a custom‐built high‐pressure cellλ=0.4246 Å. Contour plots of the HPPXRD data of Cu2(bdc)2dabco and Cu2(DB‐bdc)2dabco are shown in Figure cf. Figure S1 and S2), which includes a complete list of cell parameters, volumes, full width at half maximum (fwhm) and Rwp values as obtained from a Pawley profile fit analysis.2(bdc)2dabco a peak shift to higher 2θ angles with increasing pressure resulting from the decreasing unit cell volume is observed, see Figure p=0.175 GPa two reflections of relatively weak intensity jump from 2θ=2.63° and 2.81° to 2.75 and 2.89° with increasing intensity relative to the (100) at 2θ=2.25° and (001) at 2θ=2.53°. Furthermore, a peak broadening and intensity loss of the other reflections occurs, which is ascribed to a combination of a pressure‐induced phase transition and pressure‐induced amorphisation. The analysis of the fwhm of the reflections as a function of pressure shows an increasing fwhm from p=0.2 GPa onwards . After reaching the maximum pressure (p=0.4 GPa), the pressure was released in one step to probe for reversibility. We find partial reversibility as the cell parameters are comparable to those at p=0.2 GPa so but with a reduction of the intensity to 1/10 compared to the initial ambient measurement. This provides another indication for the combination of a phase transition and pressure‐induced amorphisation. The results are in general agreement with what has been described recently in a combined computational and Hg intrusion study on Cu2(bdc)2dabco.2(bdc)2dabco which is not seen in the Hg intrusion experiments.The HPPXRD experiments were performed in the pressure range of 2(DB‐bdc)2dabco, we find a different high‐pressure responsivity. Notably, unlike most other members of the M2(fu‐bdc)2dabco material series, Cu2(DB‐bdc)2dabco remains in its lp form after guest removal and yet no low temperature np phase has been observed.lp to np phase transition occurs at temperatures below T=100 K, not accessible by typical lab X‐ray tools; however, this makes Cu2(DB‐bdc)2dabco an interesting candidate for a high‐pressure study. Like its parent MOF Cu2(bdc)2dabco, Cu2(DB‐bdc)2dabco crystallises in a tetragonal space group, see ESI for stack plots and details of the Pawley profile fits. Up to p=0.125 GPa an isotropic compression takes place and starting from p=0.15 GPa the phase transition to the np form occurs. The phase transition occurs over a broad pressure range and is seen as a displacive phase transition according to Buerger2(bdc)2dabco the phase transition is fully reversible and the material switches back to its initial lp form with similar cell parameters compared to the initial measurement after pressure release.Looking at the high‐pressure behaviour of CuK). The bulk modulus is defined as the inverse of the compressibility and is a measure of mechanical resistance of a material towards volumetric changes under hydrostatic pressures, that is, K=-V0∂p/∂Vwith V0being the volume of the unit cell at ambient pressures and ∂p/∂Vthe change of pressure as function of volume variation. Taking into account the onset pressure of amorphisation of Cu2(bdc)2dabco and the phase transition of Cu2(DB‐bdc)2dabco, only data up to p=0.175 GPa and 0.125 GPa, respectively were fitted to a 2nd order Birch‐Murnaghan Equation of State (BM EoS) using EoSFit‐7cK(Cu2(bdc)2dabco)=14.03±0.20 GPa and K(Cu2(DB‐bdc)2dabco)=13.46±0.22 GPa. Putting these bulk moduli into the context of other reported MOFs (cf. ESI Table S4) Cu2(bdc)2dabco and Cu2(DB‐bdc)2dabco are neither particularly rigid nor soft. Interestingly, previously reported bulk moduli data for Cu2(bdc)2dabco by Wieme et al.K=11.3 GPa for Cu2(bdc)2dabco and K=10.2 GPa for Cu2(DB‐bdc)2dabco. This underlines the difficulty to obtain reliable bulk moduli for porous and rather soft materials, highlighting the large challenges related to measurement accuracies and material treatment, potential unknown contributions from defects and the difficulties in reproducing these computationally.The relatively large peak shifts in the contour plots in Figure 2(bdc)2dabco and the functionalised Cu2(DB‐bdc)2dabco were performed to obtain an atomistic picture of the ramifications of side chain functionalisation for the underlying free energy landscape. The challenge in the simulations lies in minimising the bias of the initial linker configuration on the computational outcomes within computationally accessible timescales.2(DPe‐bdc)2dabco (DPe=dipentoxy).pcu network shown in Figure S9 using the weaver code2(bdc)2dabco one single starting structure is used, whereas for Cu2(DB‐bdc)2dabco 32 individual starting structures with random phenyl orientations, that is, random torsion angles as shown in ESI Figure S9 were generated and optimized. We would like to point out that the large number of possible configurations is the reason why alternative methods to compute the free energy differences between the np and lp form based on quantum mechanical methods using internal energies plus (quasi)harmonic approximation of vibrational entropy are not readily applicable here.[Force field based molecular dynamics simulations of Cuble here.11, 57NPT and subsequent NV(aσ=0)T simulations were performed for Cu2(bdc)2dabco and Cu2(DB‐bdc)2dabco, see Figure NPT simulations, which are based on 8 (Cu2(bdc)2dabco) and 32 (Cu2(DB‐bdc)2dabco) simulations, evidence for phase transitions can already be observed; however, only from the NPT simulations it is not clear whether these simulated transition pressures originate from the underlying physics or from our simulations, and previously, premature phase transition artefacts have been observed.p(V) curves of Cu2(bdc)2dabco as obtained from NV(aσ=0)T calculations the occurrence of a phase transition is confirmed with a transition pressure of p=0.2 GPa. In contrast, a slightly different situation is observed for Cu2(DB‐bdc)2dabco where the transition pressures of each individual NPT and NV(aσ=0)T simulation are consistent with each other. For the 32 NV(aσ=0)T simulations of Cu2(DB‐bdc)2dabco, the results can be grouped into (i) a set of simulations with evidence for a structure instability which we call Group‐1  a second set of simulations with the absence of a maximum in the p(V) curve T trajectories with the experimental PXRD patterns, see ESI for details. Applying this type of rating scheme, we obtain a set of 12 simulations of the initial 32 simulations of which their PXRD patterns are in general agreement with the experiment, see Figure After setting up a methodology that allows for evaluating potential bias of the starting configuration on the computational outcomes, p(V) NV(aσ=0)T simulations of this subset we obtain a phase transition pressure of p=0.29 GPa and by including the variance of the transition pressure within individual simulations, we obtain a window of 0.19 GPa≤ptrans≤0.39 GPa which is in excellent agreement with the experiment.A more extensive discussion on the structural differences between the different subsets and their origins based on an analysis of collective variables of the different structural features can be found in the ESI. From analysing the average A was computed from p(V) by numerical integration, and the internal energy U obtained from the average of the total energy of a trajectory. For the unfunctionalised material Cu2(bdc)2dabco 2dabco is mainly driven by ΔU, and that the np form is destabilised compared to the lp form due to its unfavourable −TΔS. This behaviour is in agreement with the general rule that the crystallographically less dense structure has a higher vibrational entropy.[lp to np phase transition of Cu2(bdc)2dabco occurs rapidly and releases huge amounts of energy as the phase transition wave spreads throughout the crystallite.W=pΔV=61.6 kJ mol−1 per formula unit, and transforming it into kinetic energy of the atoms, the temperature increases by almost T=100 K. Therefore, we hypothesise that this energy is so large that it can easily overcome the energy of the coordination bonds in the MOF system, leading to amorphisation.Now we can turn our attention to the details of the underlying thermodynamics as observed from computation. The Helmholtz free energy o Figure  we see t entropy.58, 592(DB‐bdc)2dabco we observe a distinctly different behaviour 2dabco benefits from seemingly large entropic contributions compared to Cu2(bdc)2dabco with −TΔS monotonically increasing as a function of decreasing cell volume that is, increasing pressure. Given that the only difference between Cu2(DB‐bdc)2dabco and Cu2(bdc)2dabco is the alkoxy functionalisation of the bdc2− linker, we ascribe the trend in −TΔS to configurational entropy originating from side chain functionalisation. Qualitatively, reducing the available pore volume via the application of hydrostatic pressure results in a situation where side chains have a reduced number of spatial orientations, in turn reducing contributions from configurational entropy as a function of decreasing pore volume (increasing pressure). Under the assumption that entropic contributions can be divided into configurational and vibrational, and that the latter can be qualitatively described via a harmonic approximation, we can further quantify the single contributions. We optimized the lp and np forms of Cu2(bdc)2dabco and Cu2(DB‐bdc)2dabco and computed contributions from vibrational entropy at T=300 K as obtained within the harmonic approximation using phonopy.2(bdc)2dabco we obtain −TΔS=16.7 kJ mol−1, which is in agreement with the entropy penalty shown in Figure TΔS=21.1 kJ mol−1 is obtained for the functionalised material Cu2(DB‐bdc)2dabco which is significantly lower when compared to the value of approximately −TΔS=41.8 kJ mol−1 as obtained from NV(aσ=0)T simulations. Therefore, contributions from configurational entropy are calculated to −TΔS=20.7 kJ mol−1, superimposing effects from vibrational entropy and presenting an important factor in driving the entropic penalty of the functionalised MOF as the cell volume decreases. Although we here provide the exact values of vibrational and configurational entropic contributions, we would like to note that these should be regarded as qualitative. Comparing this situation to other flexible MOFs, we here show that Cu2(DB‐bdc)2dabco represents the first example in which the physicochemical properties of the flexible MOF are governed by configurational entropy. For instance, it is established that the phase transition in ZIF‐4(Zn) is based on a delicate balance between dispersion interactions and vibrational entropy, where the lp (high temperature) phase of ZIF‐4(Zn) comes with a gain in vibrational entropy driven by a softening of low frequency modes.np to lp phase transition in the well‐studied MOF MIL‐53 has been shown to be driven by vibrational entropy only.[2(DB‐bdc)2dabco is the first example in which both configurational and vibrational entropy with a weighting of approximately 1:1 exists 2dabco is a fascinating example of how small chemical modifications such as alkoxy side chains can be used to introduce configurational entropy, alter the underlying free energy landscape and therewith the macroscopic properties of a MOF. We expect that the introduction of configurational entropy via side chain modification is a general phenomenon, providing another fascinating angle of how to optimize the free energy landscape to render a rigid MOF structurally flexible.For the functionalised material Cur Figure . Firstlyopy only.10, 60s Figure . Therefo2(bdc)2dabco and Cu2(DB‐bdc)2dabco by combining state‐of‐the‐art experimentation with computation. We observe side chain dependent high‐pressure behaviour, underlining the opportunities that come with side chain functionalisation as a tool to manipulate the physicochemical properties of a MOF. The molecular dynamics simulations provide in‐depth insight into the thermodynamic factors that govern the free energy landscape, once again highlighting the important role of entropic contributions in the large structural flexibility of MOFs. In contrast to other flexible MOFs we identify configurational entropy as governing factor, drawing a clear line between the underlying thermodynamics of Cu2(DB‐bdc)2dabco and other flexible MOFs. The results suggest the use of configurational entropy as a lever for tuning the underlying free energy landscape, a factor which has been broadly overlooked as design parameter so far.In conclusion we investigated the pressure dependent responsivity of the prototypical flexible MOFs CuOur works also prove that the combination of experiment and theory can provide a detailed picture of the factors that link the microscopic arrangement of atoms with macroscopic physicochemical properties in flexible MOFs. Whilst in situ experimental methods such as solid state nuclear magnetic resonance spectroscopy and single crystal X‐ray diffraction during gas sorption experiments have greatly improved over the years,63, 64The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


Multidimensional conical intersection seam has been characterized by utilizing the dynamic resonances in the nonadiabatic transition probability experimentally observed in the predissociation of thioanisole isotopomers. 1/S2 vibronic eigenstate, providing the essential information about structure and dynamic character of the conical intersection seam projected onto the normal mode space. By modifying the nature of the normal mode space through partial or full H/D substitution of the molecule, multiple aspects of the conical intersection seam could be characterized from different viewpoints set by the adjusted normal mode space. Theoretical calculations of potential energy curves along selected normal mode displacements support the experiment.Multidimensional conical intersection seam has been characterized by utilizing the dynamic resonances in the nonadiabatic transition probability experimentally observed in the predissociation of thioanisole isotopomers. The nonadiabatic bifurcation behavior of the reactive flux into either the Herzberg type-I (electronic) or type-II  predissociation pathway is found to be strongly dependent on the quantum nature of the S The flux prepared by the 7a mode excitation is thus expected to undergo the nonadiabatic transition quite efficiently. Recently, our group has carried out the picosecond time-resolved pump-probe experiment to unravel the temporal behavior of the reactive flux associated with each of many different S1 vibronic modes of thioanisole.1/S2 conical intersection and the S1 vibrational energy is fully or partially utilized to overcome the multi-dimensional adiabatic reaction barrier. This belongs to the Herzberg type-II vibrational predissociation.0/S2 conical intersection to give the Ã state of the C6H5S˙ radical rather exclusively at the asymptotic limit. On the other hand, in the nonadiabatic channel, the reactive flux on the upper adiabatic surface above the S1/S2 conical intersection undergoes the nonadiabatic transition into the low-lying adiabatic surface most efficiently through the conical intersection, and this should belong to the Herzberg type-I electronic predissociation.0/S2 conical intersection this time with the significant probability of the nonadiabatic transition, as the reactive flux which was nonadiabatically leaked from the first S1/S2 conical intersection should have been quantum-mechanically shaped for the efficient nonadiabatic passage through the second S0/S2 conical intersection. The reaction rate on the adiabatic passage (Herzberg type-II) has been found to be much slower than that of the nonadiabatic path (Herzberg type-I). The experimental fact that the nonadiabatic transition probability strongly depends on the initial quantum state of the reactive flux indicates that the structure and energetics of the conical intersection seam could be possibly characterized in the normal mode space over the Franck–Condon region.A decade ago, our group reported the experimental result that the X[combining tilde]/Ã branching ratio (6H5S–CD3).S (S–CD3 symmetric stretching), β (CD3 bending), and an unidentified mixed mode. The dynamic resonances at multiple vibronic modes already reflect that the conical intersection is located on the multi-dimensional seam. Moreover, this suggests that the gradient vector comprising the branching plane may be described by the combination of normal modes in terms of their structural changes and associated energetics. What would be the reason why C6H5S–CD3 shows multiple dynamic resonances whereas the only 7a mode is responsible for a dynamic resonance in the undeuterated thioanisole (C6H5S–CH3)? As the electronic structures are expected to be little influenced by the H/D substitution, it is most likely that the structure and electronic property of the S1/S2 conical intersection seam would be identical for both isotopomers. The subtle differences in the nuclear displacement vectors of the normal modes of C6H5S–CH3 and C6H5S–CD3 should be then responsible for the different behaviors of dynamic resonances in these two isotopomers. In other words, the conical intersection seam associated with the sulfur–methyl bond dissociation has been projected onto different sets of normal mode coordinates for different isotopomers. In this sense, one may then be able to characterize the conical intersection seam in a number of different normal mode spaces subject to subtle changes by the partial or full H/D substitution of the molecule. As a matter of fact, our group reported that dynamic resonances at the 7a or νS mode excitation are quite robust for the particular thioanisoles of which the methyl moiety was partially deuterated .6D5S–CD3) as well as the partially deuterated in total but fully deuterated one on the benzene moiety (C6D5S–CH3). More importantly, we have carried out theoretical calculations on the first (S1) and second (S2) excited potential energy curves with respect to 7a and νS normal modes, unravelling the nature of the conical intersection seen from many different viewpoints set by different normal mode spaces of various isotopomers.In contrast to the fact that the nonadiabatic transition probability is enhanced by the excitation of the 7a mode only for the bare thioanisole, it has been found that the three more vibronic modes in addition to 7a are responsible for dynamic resonances in the partially deuterated thioanisole CH5S–CD3.66D5S–CH3; 98.5% and C6D5S–CD3; 97.17%) was purchased (MediGen Co.) and used without further purification. The gas mixture of the sample with He  was supersonically expanded through a nozzle orifice and skimmed into the differentially-pumped vacuum chamber with the backing pressure of 1–2 atmosphere at the repetition rate of 10 Hz. Detailed experimental setup had been described in previous works.via its 3Pz2A′′ (ν = 0) ← X[combining tilde]2A′′ transition of ˙CH3 (333.484 nm) or ˙CD3 (333.826 nm). The delay time between pump and probe laser pulses was fixed at 25 ns. The polarization of the pump laser pulse was parallel to a detector plate in the VMI experiment. For the two-color slow-electron velocity map imaging (SEVI)φ = 40 mm) and phosphor screen was used for both VMI and SEVI experiments. A CCD camera (IDS USB 2.0 UI-2230SE-M-GL) was employed to capture the images using the LABVIEW based acquisition program.2 dissociation dynamics.The sample . The experimental fact that the 7a mode is inactive in the R2PI spectrum, therefore, strongly suggests that the corresponding lifetime may be quite shortened. Translational energy distributions obtained by the velocity-map ion images taken from the ionization of nascent CD3 or CH3 fragment from C6D5S–CD3 and C6D5S–CH3 are shown in S modes for C6D5S–CD3, whereas it stands out only for the 7a mode for C6D5S–CH3. The X[combining tilde]/Ã product branching ratio (Γ) for C6D5S–CD3 is estimated to be 0.22 ± 0.06 or 0.34 ± 0.05 for the νS or 7a mode, respectively, compared to the value of 0.04 ± 0.03 for the zero-point energy level. This indicates that both νS and 7a mode excitations are effective in facilitating the nonadiabatic channel in the S–CD3 predissociation process while the latter seems to be more efficient. On the other hand, the X[combining tilde]/Ã product branching ratio is measured to be 0.19 ± 0.02 or 0.43 ± 0.03 at νS or 7a mode of C6D5S–CH3, respectively, compared to Γ ∼ 0.08 ± 0.04 at its S1 zero-point level. This experimental finding suggests that the 7a mode excitation of C6D5S–CH3 is much more efficient in expediting the nonadiabatic path compared to νS, 1/S2 conical intersection seam onto different normal mode platforms.Resonance-enhanced two-photon ionization (R2PI) spectra of CS or 7a, 1/S2 potential energy surfaces along the νS or 7a normal mode are so different for different isotopomers. Namely, the energetics and gradients associated with curve crossing are strongly influenced by the subtle difference of the normal mode space induced by the combinatorial way of the H/D substitution. It is remarkable that the S1/S2 curve crossing points calculated along the one-dimensional coordinate are extremely valuable for the interpretation of the experiment as below. It should be noted that the X[combining tilde]/Ã product branching ratios should linearly correlate with the bifurcation ratios of the reactive fluxes prepared near the conical intersection. Therefore, the qualitative comparison of the X[combining tilde]/Ã product branching ratios among various isotopomers based on differences in the S1/S2 potential energy surface crossings should be still reasonable. For comparison of the experiment with theoretical calculation, we will discuss the results of all four thioanisole isotopomers .Obviously, it is nontrivial to depict a multidimensional map of the conical intersection seam from the present experimental results alone although the proximity of the quantum level to the conical intersection seam is certainly correlated with the measured X[combining tilde]/Ã product branching ratio. Instead of trying to describe the conical intersection seam directly, we have carried out theoretical calculations of potential energy surfaces for various isotopomers particularly along the critical modes of ν1/S2 curve crossing along the 7a mode of C6H5S–CH3 or C6D5S–CH3, giving the X[combining tilde]/Ã product branching ratio of 0.416 ± 0.0333 or 0.43 ± 0.03, respectively  in 1 state survives long enough during the laser ionization process by the nanosecond laser pulse. Overall, the S1/S2 crossing point along the 7a mode is found to be located around 1200 cm–1 above the S1 minimum energy point for C6H5S–CH3 and C6D5S–CH3, –1 for the 7a mode, this energetic position of the curve crossing point is quite close to the first quanta of the 7a mode, suggesting that the conical intersection is likely to be easily accessed by the 7a mode excitation of those molecules. For all other isotopomers, the S1/S2 curve crossing points along the 7a mode are predicted to be located at ∼2000 cm–1 or less above the S1 minimum energy points, 6H5S–CD3 or C6D5S–CD3, respectively, reflects that the corresponding vibronic excitation prepares the reactive flux at the conical intersection seam along which the S–CD3 bond elongation should be effectively coupled. At the present time, it is less straightforward to quantify the structural and energetic characteristics from the comparison of the experimentally measured X[combining tilde]/Ã branching ratio with the calculated potential energy surfaces.The most dramatic shape of the conical intersection, for example, could be inferred from the calculated Sectively . In thesS mode excitation of C6H5S–CD3 giving the X[combining tilde]/Ã product branching ratio of 0.33 ± 0.04, whereas the νS mode excitation is not so effective in facilitating the nonadiabatic transition for C6H5S–CH3 to give Γ ∼ 0.165 ± 0.025.1/S2 crossing point is calculated to be more accessible by the one quantum of νS for C6H5S–CD3, 6H5S–CH3, 1/S2 crossing point along νS is located energetically much higher for C6H5S–CH3 compared to that of C6H5S–CD3 with respect to their S1 minima. Moreover, the S1/S2 crossing point of C6H5S–CH3 is quite remote from the S1 minimum energy position in terms of the nuclear displacement compared to that of C6H5S–CD3. It is indeed remarkable that the simple potential curve calculation could explain the experimental finding very well at least qualitatively. The effect of deuteration of the methyl moiety on the reaction nonadiabaticity seems to be similarly applicable to isotopomers of C6D5S–CH3 and C6D5S–CD3 at the νS mode excitation, giving the X[combining tilde]/Ã branching ratio of 0.19 ± 0.02 and 0.22 ± 0.06, respectively. The S1/S2 curve crossing points of two isotopomers along νS support the experiment, as the crossing point of C6D5S–CD3 is predicted to be located lower than that of C6D5S–CH3 with respect to the S1 minima of those isotopomers, S mode excitations of C6H5S–CH3 and C6D5S–CH3 are hardly associated with dynamic resonances in the proximity of the conical intersection seam, and thus the explanation of those experimental results based on the potential energy curve crossings seems to be less meaningful at the present time. The dynamic role of the pseudo-conical intersection previously predicted to be existed at the C–C–S–CH3 dihedral angle of 90°1 internal energy region where the out-of-plane modes are likely to be more activated.One of the most puzzling experimental observations has been the fact that the X[combining tilde]/Ã product branching ratio shows the strong resonance-like enhancement at the ν1/S2 conical intersection seam onto the normal mode space could be inferred from the nature of the S1/S2 vibronic modes responsible for dynamic resonances showing the large enhancement of the nonadiabatic transition probability. The X[combining tilde]/Ã product branching ratios estimated from the translational energy distributions of the nascent methyl fragments measured for various isotopomers of thioanisole are particularly useful as the proximity of the initial reactive flux to the conical intersection seam could be judged from different viewpoints set by the different normal mode space. Accordingly, the nonadiabatic bifurcation dynamics of the reactive flux into either the Herzberg type-I (electronic) or type-II  predissociation pathway has been disentangled and the structure and dynamic role of the conical intersection seam could be inspected from various points of view. With the aid of simple calculations of potential energy curves along selected normal mode displacements, we report the first cornerstone for the otherwise formidable characterization of the conical intersection seam in polyatomic molecular system.Quantum mechanical characterization of the multidimensional conical intersection seam has been demonstrated to be feasible in predissociation reaction of the excited thioanisole molecule. The projection of the SThere are no conflicts to declare.Supplementary informationClick here for additional data file.


Herein, intrinsic long‐range ferromagnetic order in 2D Ta3FeS6 is reported, which exhibits ultrahigh stability under the atmospheric environment. The intrinsic ferromagnetism of few‐layer Ta3FeS6 is revealed by polar magneto‐optical Kerr effect measurement, which exhibits giant MOKE response and has Curie temperature of ≈80 K. More importantly, few‐layer Ta3FeS6 nanosheet exhibits excellent air stability and its ferromagnetism remains unchanged after 4 months of aging under the atmosphere. This work enriches the family of 2D ferromagnetic materials, which will facilitate the research progress of spintronics.2D ferromagnetic materials provide an important platform for the fundamental magnetic research at atomic‐layer thickness which has great prospects for next‐generation spintronic devices. However, the currently discovered 2D ferromagnetic materials  of ≈80 K. Even more important, Ta3FeS6 exhibits excellent air stability and its ferromagnetism remains unchanged after 4 months of aging in air, indicating its potential application in spintronic devices under atmospheric environment.Ultrastable intrinsic ferromagnetism in two‐dimensional Ta Recently, 2D ferromagnetism was discovered in chromium trihalides , Cr2Ge2Te6, and Fe3GeTe2. Magnetic CrI3 exhibits electrical‐field tunable 2D magnetism, large interlayer tunneling magnetoresistance, and helical luminescence properties. However, most of the found 2D magnetic materials have poor stability in the atmospheric environment, i.e., CrI3 nanosheet degrades in the air in 15 min and the ferromagnetism of Fe3GeTe2 nanosheet vanishes under the atmosphere for a few hours, which hinders the scientific research of intrinsic 2D ferromagnetism and the practical applications. Therefore, it is significant to explore high air‐stability 2D ferromagnetic materials for the research of 2D intrinsic ferromagnetism and the potential application of spintronics.2D materials have generated great research interest due to their atomic flat interface structure, unique electronic structure , and a wide range of applications .5, 63FeS6 was demonstrated with excellent air stability and strong magnetocrystalline anisotropy, which would be an ideal platform for investigating the 2D magnetism and spin devices. 2D ferromagnetic Ta3FeS6 nanosheets with different thickness were successfully prepared by the combination of crystal growth technique and Au‐assisted mechanical exfoliation strategy. Polar MOKE measurement demonstrated the robust 2D ferromagnetism of Ta3FeS6 with a giant MOKE rotation angle of 48 mrad. Further, temperature‐dependent MOKE measurement demonstrated that the Curie temperature of Ta3FeS6 nanosheet was 80 K. What's more, Ta3FeS6 nanosheet had excellent air stability and its ferromagnetism remained unchanged after 4 months aging. The air‐stable 2D ferromagnetism of Ta3FeS6 is an ideal platform for the investigation of spin devices and the construction of van der Waals magnetic heterostructures.In this study, 2D intrinsic ferromagnetism in Ta3FeS6 was isolated from high‐crystalline single crystal via mechanical exfoliation. The Ta3FeS6 single crystal was grown by the chemical vapor transport (CVT) method, as illustrated in Figure 3FeS6 belongs to the space group of P6322 (No. 182) and has hexagonal layered lattice structure with noncentrosymmetric and chiral feature, which is alternately stacked by Fe atoms and H‐phase TaS2 layers, as illustrated in Figure  The crystal can also be denoted as Fe1/3TaS2 due to the 3a×3a Fe atoms superlattice. The interlayer spacing of neighboring TaS2 is 0.614 nm and is larger than that of 2H‐TaS2 (0.605 nm) because of the intercalation induced lattice expansion. Besides, the Fe atoms in the adjacent layers have a complementary occupation and constitute honeycomb Fe lattice along the c axis of Ta3FeS6 crystal  was employed to confirm the phase structure of Ta3FeS6 crystal. As shown in Figure 3FeS6 crystal (JCPDS No. 22‐0360) and further demonstrated its layered features with the high‐crystalline quality.2D Tan Figure 22 Thel Figure .[23]2 layers, the isolating of Ta3FeS6 nanosheets in experiment is still a great challenge. In recent years, metal‐assistant isolation technique was developed and had been proved to be efficient for large‐area preparation of various layered materials due to the chemically enhanced adhesion force between evaporated metal film and surface of materials. Hence, the Au‐assisted mechanical exfoliation method was introduced to overcome the strong interlayer interaction and prepare atomic‐thin Ta3FeS6 nanosheet. The exfoliation process was illustrated in the Experimental Section and Figure S1 (Supporting Information). Figure 3FeS6 nanosheet on 200 nm Au film and the few‐layer regions can be resolved by the optical contrast. The topography characterization was performed by an atomic force microscope (AFM) to identify the thickness of the Ta3FeS6 nanosheet. AFM image demonstrated that the cleaved nanosheet had bi‐ to eight‐layer thickness with distinct layer steps of ≈0.6 nm matching well with the interlayer distance of Ta3FeS6 crystal.However, due to the ionic bonding effect between the Fe atoms and the TaShallenge.25 In 3FeS6 nanosheet, transmission electron microscope (TEM) measurement was performed on an exfoliated nanosheet. Figure 3FeS6 nanosheet on a carbon grid and the transparent feature indicated its ultrathin character. High‐resolution TEM (HRTEM) image  analysis was carried out. The result showed that the as‐synthesized crystal contains Ta, Fe, and S elements with the ratio of approximate 3:1:6, as shown in Figure 3FeS6, as shown in Figure 2 frame. More intriguing, another concentric set of hexagonal spots was observed at (1/3 1/3 0) reciprocal positions, resulting from an ordered Fe superlattice of Ta3FeS6. All of the analysis above agreed well with the structural feature of Ta3FeS6 crystal. Additionally, the phase homogeneity of the Ta3FeS6 nanosheet was revealed by SAED. The SAED patterns , and P5 was the transverse optical vibration mode (A1g mode). The Raman spectra at 80 and 300 K maintained a constant shape indicating that no crystal structure transition occured during the temperature‐changing process. To visually present the evolution of temperature‐dependent Raman spectra, the 2D contour map ranging from 80 K to 300 K was plotted, as shown in Figure 2, P4, and P5 exhibited obvious red‐shift (soften behavior). To analysis the soften behavior quantitatively, the relative peak position offset  were extracted and shown in Figure 5 peak showed linear behavior along with temperature which can be fitted by the Lorentzian functions, as shown in Equation  × 10−2 cm−1 K−1, which was close to the reported first‐order temperature coefficient of other TMDs, such as MoS2  and WS2 . As demonstrated in previous works, the first‐order temperature coefficient is relative to the bonding‐type and structural configuration of materials. The similar χ value compared with MoS2 and WS2 indicated the layered nature in Ta3FeS6 crystal.Raman spectroscopy was further employed to investigate the spin dependence of Raman scattering of Taansition.32Fig2 peak offset occured at around 145 K, as indicated in Figure χ of the low‐temperature region (80–145 K) was −(8.7 ± 1.7) × 10−2 cm−1 K−1, which was 7.9 times larger than that at high‐temperature region (−(1.1 ± 0.2) × 10−2 cm−1 K−1, 145–300 K). Previous reports indicate that some of Raman peak positions or intensities change with decreasing temperature through the Curie temperature. The prominent difference of Raman spectra between paramagnetic phase and ferromagnetic (or antiferromagnetic) phase can be summarized as follows: 1) New Raman modes due to the magnetic ordering (two‐magnon scattering); 2) Quenching of quasi‐elastic scattering from magnetic fluctuations; 3) Magnetic ordering induced Raman peak offset (frequency difference). Since the Curie temperature of Ta3FeS6 was ≈145 K, it can be deduced that the abnormal Raman offset was attributed to the ferromagnetic‐paramagnetic phase transition of Ta3FeS6, which was also observed in NiPS3. Thus from this perspective, the P2 Raman peak can be used as an indicator of the magnetic ordering for Ta3FeS6.More intriguing, a turning point of Pperature.41 TheFigure 3FeS6. To obtain the out‐of‐plane ferromagnetism, a normal‐irradiated laser was used to detect the MOKE signals of samples. The schematic for the optical setup of polar MOKE is illustrated in Figure S3 (Supporting Information). Figure θK as a function of the magnetic field at different laser wavelengths for the Ta3FeS6 nanosheet . All the experiments were measured at 10 K with an incident light power of 3 µW. To rule out the possible contribution of Kerr signal by Au residues, the MOKE measurement was performed on bare Au film under the same condition and no obvious MOKE signal was observed . Significantly, the magnetic hysteresis loop with a sharp switching edge demonstrated that the Ta3FeS6 nanosheet was ferromagnetic ordering with strong out‐of‐plane spin polarization. Previous theoretical calculation demonstrates that the local moments at Fe atoms induces a magnetic polarization of the Ta 5dZ2 band, and the direction of polarization was parallel arranged along the c axis in the Ta3FeS6 structure, which agreed well with our experiment result. Furthermore, the MOKE signals displayed obvious dependence with the incident wavelength that some signals showed positive and others showed negative. The values of Kerr rotation θK and coercive field HC were extracted as functions of wavelength, as shown in Figure 3FeS6. However, the Kerr rotation θK exhibited strong wavelength dependence and changed its sign at ≈540 nm. The Kerr signal is proportional to the difference of spin‐up and spin‐down transition rate, which equals to joint density of states (JDOS). It is worth noting that the difference of JDOS between the spin‐up and spin‐down bands has photon energy dependence. When JDOS of the spin‐up band is larger than that of the spin‐down band, the Kerr signal shows positive and vice versa. Hence, the sign of Kerr signal changes at ≈540 nm may attribute to the sign change of the difference of JDOS.Polar MOKE measurements have the advantage of high sensitivity, high spatial resolution, and non‐invasive detection, which are propitious to the research of 2D magnetic materials.43Fig3FeS6 nanosheet. Figure θK as a function of the magnetic field at several temperatures. With the increase of temperature, the hysteresis loop finally shrank away, indicating that the ferromagnetic phase was transforming into the paramagnetic phase. Figure θK and the coercive field from Figure 3FeS6 nanosheet was zero, which implied that the Curie temperature TC was around 80 K. Compared to the bulk sample , the suppressed ferromagnetism of Ta3FeS6 nanosheet was due to the thermal fluctuations in the 2D system. Layer‐dependent ferromagnetism of Ta3FeS6 has been carried out, but it was hard to figure out the layer‐dependent phase diagram owing to the inhomogeneity of Fe atoms in the Ta3FeS6 crystal. The Fe content of Ta3FeS6 changed after the mechanical exfoliation process, as demonstrated by the EDS analysis . Besides, the Curie temperature was sensitive to the Fe content in Ta3FeS6 crystal . When there were more Fe vacancies in the superlattice, the Curie temperature of Ta3FeS6 crystal increased from 80 K to 145 K . The Fe vacancies arose from the crystal growth process. During the crystal growth process, the Fe atoms diffused to the TaS2 frame. Affected by the growth condition, such as the non‐homogeneous temperature distribution, the inhomogeneous diffusion of Fe atoms resulted in the fluctuation of Fe concentration in different layers of Ta3FeS6 crystal.Temperature‐dependent polar MOKE measurements were performed to determine the Curie temperature of the Ta Here, aging measurement was performed for researching the stability of 2D Ta3FeS6. After 4 months aging in air, the coercive field and Kerr rotation θK of Ta3FeS6 nanosheet remained unchanged, as shown in Figure 3FeS6 nanosheet was superior to recent reported 2D ferromagnetic materials, like CrI3, Cr2Ge2Te6, and Fe3GeTe2. The high stability of 2D Ta3FeS6 may attribute to the larger electronegativity of sulfur and its chemical inertness to H2O and O2, suggesting its potential application in spintronic devices under the atmospheric environment.A significant obstacle for the research and application of 2D ferromagnetic materials is their air stability.15 Her3FeS6 and investigated the ultra‐stable 2D ferromagnetism. Giant MOKE response was demonstrated in Ta3FeS6 nanosheet with Kerr rotation angle of 48 mrad and Curie temperature of ≈80 K. Most importantly, the Ta3FeS6 nanosheet exhibits outstanding air stability and its MOKE signal remains unchanged after 4 months of aging in air. We envision that the long‐life 2D ferromagnetic ordering of Ta3FeS6 will provide more opportunities for field‐effect modulation of magnetism and Curie temperature, magnetic doping engineering, and fabrication of magnetic tunneling junction working at the atmospheric environment.In summary, we have successfully prepared 2D Ta3FeS6 was prepared by mechanical exfoliation of Ta3FeS6 single crystal. Ta3FeS6 single crystals were synthesized by the I2‐assisted CVT method. High‐pure Ta , Fe , and S  was mixed (molar ratio of 1:0.4:2) and then sealed under vacuum in a quartz tube with the addition of 200 mg I2 as the transport agent. Whereafter, the ampoule was placed in a two‐temperature zone horizontal tube furnace with the hot end at 1273 K and cold end at 1173 K in 10 h and kept for a week. Ta3FeS6 nanosheets were obtained by Au‐assisted as illustrated in Figure S1 (Supporting Information). First, single crystal Ta3FeS6 is placed on adhesive tape. Then Au film (200 nm) is evaporated onto the single crystal by thermal deposition. The Au atoms bond with the surface S atoms of the crystal. Next, the Au film was picked up by another tape along with some Ta3FeS6 flakes cohered on the surface. Finally, it's necessary to reduce the thickness of Ta3FeS6 flakes on Au film by mechanical exfoliation for several times.2D TaT method.52 Hig3FeS6 crystals were characterized by an OM , XRD , a Raman spectrometer  equipped with a cryo console, and an AFM . The TEM, SAED, and EDS were performed in a field emission TEM . The Au film was deposition in a high‐vacuum deposition system . MOKE measurements were carried out by using a home‐made microscopic polar MOKE system. The samples were placed in a superconducting magnet with a temperature range from 4.2 to 300 K and an out‐of‐plane magnetic field up to 5 T. The wavelength of laser source range from 500 to 900 nm with an excitation power of 3 µW. Incident light is normally incident on the sample surface through an aspheric lens. Mechanical chopper can change the incident light intensity and photoelastic modulator (PEM) can modulate the polarization of incident light. The combination of mechanical chopper and PEM can extract the MOKE signal of the sample. The detail of the optical setup is shown in Figure S3 (Supporting Information).TaThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


In this work, we synthesized and investigated the crystals of a novel mixed‐valent iron oxide with an unconventional Fe5O6 stoichiometry. Near 275 K, Fe5O6 undergoes a Verwey‐type charge‐ordering transition that is concurrent with a dimerization in the iron chains and a following formation of new Fe−Fe chemical bonds. This unique feature highlights Fe5O6 as a promising candidate for the use in innovative applications. We established that the minimal Fe−Fe distance in the octahedral chains is a key parameter that determines the type and temperature of charge ordering. This model provides new insights into charge‐ordering phenomena in transition‐metal oxides in general.Functional oxides whose physicochemical properties may be reversibly changed at standard conditions are potential candidates for the use in next‐generation nanoelectronic devices. To date, vanadium dioxide (VO Charge and order: Crystals of novel mixed‐valent iron oxide with an unconventional Fe5O6 stoichiometry were synthesized and investigated. Near 275 K, Fe5O6 demonstrates a Verwey‐type charge‐ordering transition that is concurrent with a pronounced dimerization and the formation of new Fe−Fe chemical bonds. This unique feature highlights Fe5O6 as a highly promising candidate for use in innovative applications. These novel mixed‐valence iron oxides are of fundamental importance, but their physicochemical properties and technological potential have yet to be investigated.Strongly correlated oxides, with their unusual optoelectronic and magnetic properties related to interacting electrical charges, are of considerable interest from both fundamental and applied perspectives.5O6 stoichiometry,5O6 undergoes an unusual phase transition, with charge ordering realized through the formation of novel Fe−Fe chemical bonds between two adjacent ions in the linear octahedral chains. This leads to a pronounced dimerization within these chains and to significant changes in the electrical resistivity and other physical properties. Even though the phase transition in Fe5O6 is not as abrupt as the metal–insulator transition in VO2,5O6 whose orientation coincides with the dimerization direction. From the more fundamental perspective, we establish that the minimal Fe−Fe distance in the octahedral chains is the key parameter that determines the type and temperature of the charge‐ordering transition. This model provides new insights into charge‐ordering phenomena in both iron oxides and transition‐metal oxides in general.Here, we synthesize crystals of a novel iron oxide with the unconventional Fe5O6 samples under moderate HP‐HT conditions using a multi‐anvil press at Bayerisches Geoinstitut 5O6 samples adopted the orthorhombic Cmcm structure (No. 63), where the iron cations fill the linear chains of both the octahedra  and the trigonal prisms .2+, likewise, the two inequivalent octahedral sites are filled with mixed Fe2+/Fe3+ ions.We synthesized the Fe5O6 single crystals . The Fe−Fe distances in the octahedral chains in this phase displayed a pronounced separation into pairs, resulting in the formation of dimers , which reveals a shallow but well‐defined minimum around T≈270 K  curves observed upon cooling below 270 K suggests the formation of local magnetic moments following charge localization on the iron dimers. Non‐linear inverse susceptibility indicates deviations from the Curie–Weiss behavior in this temperature range  curves marks a purely antiferromagnetic transition with the Néel temperature of TN=100 K in the applied field of μ0 H=1 T  curves below TN in higher fields. The isothermal magnetization measurements showed a series of metamagnetic transitions around TN for 5 T<μ0 H<7 T  of several iron oxides vs. the shortest Fe−Fe distances, d(Fe−Fe)oct, in the edge‐shared linear octahedral chains. This plot indicates a crossover near d(Fe−Fe)oct≈2.91 Å separating the regions of short and long Fe−Fe distances, optimal for the formation of dimers and trimers, respectively  in the prisms and nearly identically charged octahedra (Fe2.67+ on average)—demonstrate different types of charge order. Upon the charge ordering in the least dense CaFe3O5 with d(Fe−Fe)oct=3.021 Å or 3.014 Å, the iron atoms shift by only ≈0.01 Å and form loosely packed iron trimers,3O5 with d(Fe−Fe)oct=2.914 Å lies near the crossover point and exhibits a conventional Fe2+/Fe3+ charge separation.4O5, which showed a competition between the trimeric and dimeric types of charge ordering, leading to a complex incommensurately modulated structure of the charge‐ordered phase.4O5 could be realized due to a compromise between the high Fe2.67+ average charge in the octahedra, which is optimal for the formation of trimers (Fe3+−Fe2+−Fe3+),d(Fe−Fe)oct=2.891 Å distance that is more suitable for the formation of dimers oct=2.877 Å distance, optimally charged octahedral ions (Fe2.5+ on average), and with the high TCO=275 K temperature further enhanced under pressure fits well to the trend found for Fe4O5 Fe4O6,The coupling of the charge ordering and dimerization we observed in FeThe authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.