Herein, dual rotating liquid film reactors (RLFRs) and lime are proposed to construct a recyclable, ultra‐fast, and value‐added desulfation method. Parameter optimization and kinetic calculations prove that the above reactions are controlled by internal diffusion, revealing that RLFR promotes the mass transfer and reaction rate. The new process greatly shortens the desulfation time of lead paste from 40 min to 10 s with a high desulfation rate of 99.7%, and the sulfation time of lime from 30 min to 30 s with a sulfation rate of 98.6% with a net profit of 55.99 ¥/ton by cost accounting. Moreover, ten batches of continuous scale‐up experiments demonstrate the stability of processes, the desulfation and sulfation rates are kept at 99.7% and 98.2%, which greatly reduces the emissions of waste desulfate liquor. This work provides a new universal strategy for a sustainable, low‐cost, and clean desulfation method of waste resources to achieve technical and economic feasibility.The new low‐cost clean pre‐desulfation technology is very important in pyrometallurgy and hydrometallurgy. However, traditional reactors have low space‐time yield and desulfation rate, resulting in high energy consumption and SO This work proposes a new ultra‐fast recyclable intensification desulfation process of lead paste and a mother liquor regenerated by lime via dual rotating liquid film reactors to grind the surface of particles and enhance the mass transfer process, which achieves the desulfation rate and conversion rate of 99.7% and 98.6% in 10 and 30 s, respectively. Lead is an important nonferrous heavy metal widely used in lead‐acid batteries (LABs), chemical equipment, nuclear industries, and cable communication. The emission of heavy metal lead in LABs causes severe damage to human health in the brain and nervous system. With the rapid sustainable development of the energy storage industry in emerging economies, 8.3 million tons of refined lead are used to manufacture the production of LABs, which accounts for over 80% of the total global refined lead production in 2022. From the point of view of resource consumption or the environment, a large number of spent LABs for effective and rational recycling are the focus of attention to achieve the sustainable development of lead resources and reduce environmental pollution. Spent lead paste, formed by the active material after long‐term charge and discharge, is the high‐quality lead‐containing resource recovered from spent LABs, which mainly contains PbSO4 (50–60%), PbO2 (15–35%), PbO (5–10%), Pb (2–5%), and other impurities components (1%) such as Fe, Ba, Al, Sn, and Sb. Among components of spent lead paste, PbSO4 is the main leady compound and is most difficult to be dissolved in water and acid. It is also thermodynamic stable, even under a high temperature of 1300 °C. The decomposition of PbSO4 phases leads to the emission of SO2 and lead particulates. The desulfation process in pyrometallurgy generates emissions of SO2 and lead dust and requires a lot of energy. The desulfation process in hydrometallurgy has technical bottlenecks such as expensive desulfurizers, low leaching rate, long reaction time, and high disposal cost of liquid waste treatment. Therefore, from the point of view of resource consumption or environmental protection, the pursuit of a clean, energy‐saving, value‐added, highly efficient, and ultra‐fast pre‐desulfation process of lead paste is very important in the smelting process and hydrometallurgy.Power batteries are considered the top priority for new energy vehicles to achieve the carbon neutrality goal.1, 2 Most of the reported desulfurizers, such as Na2CO3, K2CO3, CaCO3, (NH4)2CO3, NH4HCO3, and NaOH, can convert PbSO4 into PbCO3 or Pb(OH)2 with a dissatisfactory desulfation rate of 95–99% and a long desulfation time of 10–120 min, which is because the solubility of insoluble PbCO3 is much smaller than that of PbSO4 in the spontaneous solid–liquid phase reaction, resulting in a sluggish kinetics process, reaction time, and conversion rate. Recently, Yang et al. used NaOH as a desulfurizer to desulfurize in the hydrothermal reactor with a high desulfation rate of 99.46% for 1 h. Liu et al. compared the desulfation effect of Na2CO3, (NH4)2CO3, and NH4HCO3 on the spent lead paste, and found that the type of desulfurizer displayed a little effect on desulfation time and rate, and selected 1 h as the optimal operation time. Expensive desulfurizers of NaOH and Na2CO3 with subsequent evaporation cost of sulfate mother liquor, and slow desulfation speed greatly increase the cost of desulfation, hindering the industrialization application of pre‐desulfation technology. In addition, the Junqing Pan group reported a lime ex situ nucleation + ammonia‐CO2 cycle closed loop desulfation process to reduce the cost of desulfurizer. The ex situ nucleation process greatly improves the utilization rate of the lime conversion process, achieving a desulfation rate of 99.4% within 60 min. In terms of reactor modification to improve the reaction speed, Zhang et al. reported the “surface update” process to significantly short the desulfation time by quickly removing the desulfation product from the surface of PbSO4 particles with the help of mechanical force, which achieved an extremely low PbSO4 residue of 0.5% and ultra‐fast desulfation time of 20 min compared to the traditional process (2% and 60 min). Ma et al. proposed a forced surface renewal mechanism for adaptive surface grinding based on magnetic levitation, which greatly improved reaction efficiency and mass transfer speed, resulting in an increased desulfation rate of 97% at 3 h. Ning et al. used the rotating packed bed as the reactor for enhanced desulfation of spent lead paste, achieving a high desulfation rate of 97.2% in a discontinuous desulfation process.The reported new desulfation technologies include new desulfation agents, desulfation equipment, and related technologies.15 MosThe major reasons for the low desulfation rate are the following problems: 1) Most stirred reactors own limited space and low shear force, leading to a low conversion rate and increasing energy consumption due to the unreacted raw materials to be covered by the generated product coating layer. 2) Most stirred reactors have long reaction times and low space‐time yield due to low kinetic rate. 3) Most of the stirred reactors possess a discontinuous reaction process, resulting in low production efficiency and inconsistent products of different batches. Therefore, it is urgent to design and adopt a new highly efficient intensification reactor and inexpensive desulfation agents to construct a fast and continuous desulfation process to promote the future industrial application of pre‐desulfation.Scheme4 particles and strengthens the mass transfer effect through the collision grinding effect of liquids and particles generated by the directional high‐speed rotation process in the reaction process. The new processes of desulfation of PbSO4 and sulfation of Ca(OH)2 can be instantly completed in 10 and 30 s, respectively, which significantly enhance reaction speeds, conversion efficiencies, and space‐time yield, achieving a directional, continuous, and sustainable technology, guiding a clean revolutionary improvement on the treatment of lead industry and recycling of the other related solid waste resources.Herein, we propose a new ultra‐fast recyclable desulfation process of spent lead paste via the rotating liquid film reactor (RLFR) to grind the surface of particles and enhance the mass transfer process Scheme1. Expe22.1 Compared with the low reaction efficiency, long reaction time, and discontinuous reaction caused by the reaction products coated on the unreacted raw materials in the traditional reactor, 1) the RLFR possesses the tremendous shear force generated by the rotor‐stator system, which can make the prepared materials have the characteristics of small average particle size and narrow particle size distribution, greatly enhancing the microscopic mixing and mass transfer process. 2) The RLFR with the small confinement space and high rotation speed, makes raw material and liquid rapidly pass through the gap between the high‐speed rotor and stator under the action of extrusion, which causes the continuously exposed new surfaces to contact with the desulfation agent, thus greatly shortening the desulfation time and improving the reaction efficiency. Thus, the new device instantly generates a large grinding force and enough transformation chance among material particles, leading to the forced renewal material surface and the instantaneous accomplishment of the reaction.In this case, the RLFR is proposed as a new intensification device to rapidly collide raw materials efficiently by instantaneous solid–liquid surface renewal with a greatly increased mass transfer rate to break through the diffusion‐controlled chemical reaction process. During the working process of the RLFR, the gap between the rotor and the stator is fine‐tuned, and the materials collide and grind in high‐speed movement through the gap, thereby the materials can achieve four purposes of mixing, dispersing, pulverizing, and reacting in the RLFR, and finally, fine productions are obtained.31 Com2.24)2CO3 and PbSO4 and the sulfation reaction of Ca(OH)2 and (NH4)2SO4 are spontaneous processes with the reduced Gibbs free energy value (△rGm) according to Lange's Handbook of Chemistry. Based on the solubility product (Ksp) values of PbSO4 (Ksp = 1.6 × 10−8), PbCO3 (Ksp = 7.4 × 10−14), Ca(OH)2 (Ksp = 5.5 × 10−6), and CaSO4 (Ksp = 9.1 × 10−6), it is known that these two solid–liquid phase reactions are ways in which one insoluble compound generates another more insoluble compound. Due to the low solubility of these two solid–liquid phase reactions, resulting in a sluggish kinetics process, it is a great primary challenge to accelerate the actual reaction speed and conversion rate during the reaction. In this paper, Figure4)2CO3 and the sulfidation process of (NH4)2SO4 mother liquor and Ca(OH)2 by the intensification process of RLFR, which improves the inherent shortcomings of traditional reactors, revealing the versatility for different types of reactions.Thermodynamically, the desulfation reaction (1) of 2CO3 solution is instantly ground and pulverized through multiple functions of high‐speed shearing, squeezing, and impacting under vertical gravity, which promotes the forced renewed of the surface of spent lead paste and continuously exposes a new surface to contact with the desulfurizer, greatly shortening the desulfation time and improving the reaction efficiency.Figure 4)2CO3 to PbSO4, (NH4)2CO3 concentrations, the distance between the stator and the rotor, and reaction temperatures on the desulfation rate. With the increase of molar ratios of PbSO4 to (NH4)2CO3 from 1:1 to 1:1.7, the desulfation rate increases from 91.6% to 99.6% at 10 s 2CO3 to 1:2.2, delivering that an appropriate molar ratio can accelerate the desulfation process and the reaction rate. Figure 4)2CO3 on the desulfation rate. It confirms that the increase in concentration and viscosity of (NH4)2CO3 solutions, resulted in a reduction of the desulfation rate. Figure 4)2CO3 at 60 °C. Benefitting from its high‐speed grinding and enhanced liquid dispersion of RLFR, it achieves up to 240 times desulfation speed than traditional reactors at room temperature, demonstrating a remarkable process intensification effect. So the optimal desulfation conditions of this process are as below: a molar ratio of PbSO4/(NH4)2CO3 is 1:1.7, the (NH4)2CO3 concentration is 0.5 mol L−1, the reaction time is 10 s, a distance between the stator and rotor is 0.5 µm, and the desulfation efficiency of spent lead paste is as high as 99.7% under room temperature.To deeply reveal the intensification desulfation process via the RLFR, we have systematically studied the effects of reaction times, different molar ratios of 2CO3 is a solid–liquid phase reaction, with most of the reaction occurring at the interface. It is assumed that the desulfation reaction of PbSO4 reacts from the outside to the inside. Firstly, when the reaction time progresses, reaction (1) continuously shrinks to the internal solid and generates a new PbCO3 product layer, accompanied by gradual thickening and encapsulation of the PbCO3 product on the surface of solid raw material, leading to a slowing down of the reaction and a shrinking core of unreacted raw material solid. This reaction can be characterized by the shrinking core model of the PbSO4 desulfation .If the reaction process is controlled by internal diffusion, the reaction kinetic equation can be represented as30(4K isT is the thermodynamic temperature (K). K is the reaction rate constant at temperature; A is the pre‐exponential factor; R is the thermodynamic gas constant (8.314 J mol−1 K−1). Ea is the apparent activation energy (kJ mol−1).According to Arrhenius formula, the relationship between temperature and reaction rate constant K is37 between 1−(2/3)x−(1−x)2/3 and reaction time (t) is significantly better than that between 1−(1−x)1/3 and t under different reaction temperatures, further indicating that the reaction between PbSO4 and (NH4)2CO3 is dominated by the internal diffusion. Thus, the KD of Equation  is 5.82 kJ mol−1 by analyzing the linear relationship between 1/T and InK. Because the Ea value is not in the range of 30–85 kJ mol−1, it is inferred that the process is not controlled by surface chemical reactions but by internal diffusion.Figure T) and K can be obtained from the Arrhenius equation, as shown in the following Equation KT2Ea = 4.20 kJ mol−1, T1 = 293.15 K, and T2 = 303.15 K are substituted into Equation K30T1 = 313.15 K and T2 = 323.15 K are substituted into Equation (When Equation (8)K32K values increase slowly with the increase of T. It is the low Ea value that leads the reaction to be insensitive to temperature changes. Therefore, the introduction of new catalysts can reduce the Ea values of this desulfation reaction and the influence of temperature on the desulfation rate, which will become the focus of follow‐up research.According to the above calculation results, the 4)2CO3 on the desulfation reaction rate is further evaluated as shown in Figure R from the fitting curves exhibits that the desulfation process is still dominated by internal diffusion while changing (NH4)2CO3 concentrations. The slope of 1.03 in the linear regression relationship between lnK and lnC is the apparent reaction order of the desulfation reaction, indicating that the reaction rate is proportional to the concentration to the power of 1.03 2CO3 to PbSO4 is 1.7:1, the desulfation rate reaches 99.7% after 10 s and the consumption of (NH4)2CO3 and PbSO4 0.2991 mol involved in the reaction (1). Based on the above kinetic process analysis, the desulfation process of spent lead paste is still controlled by internal diffusion. Therefore, the intensification effect of shearing, mechanical grinding, and ultrasonic crushing is the key to improving the desulfation rate of lead paste. It explains the great contribution of RLFR to enhance the conversion speed and desulfation rate of lead paste and also indicates that the desulfation conversion rate of lead paste is not sensitive to changes in reaction temperature and (NH4)2CO3 concentration.When the optimal molar ratio of (NH2.2.32 and (NH4)2SO4 regenerates NH3·H2O solution and CaSO4, it found that formed CaSO4 was easily covered on the unreacted Ca(OH)2, resulting in the slow reaction rate and a low conversion rate. In this work, RLFR is introduced to enhance the conversion efficiency of Ca(OH)2 by systematically studying the effects of reaction times, different molar ratios of Ca(OH)2 to (NH4)2SO4, Ca(OH)2 concentrations, the distance between the stator and the rotor, and reaction temperatures on the sulfation rate. With the increase of molar ratios of Ca(OH)2 to (NH4)2SO4 from 1:1 to 1:1.8, the sulfation rate increases from 97.8% to 98.3% at 30 s 2SO4 can achieve a conversion rate of 98–99% within 3 min, which corroborates that the molar ratio of raw materials has little effect on the conversion rate of Ca(OH)2. The remaining 1% is impurities from hydrated lime or tiny core particles. Figure 2 on the sulfation rate, which suggests that the increase in concentration and viscosity of (NH4)2CO3 solutions results in a reduction of the sulfation rate.The reaction of Ca(OH)ion rate.25 In 2 from Figure 3·H2O solution and increase the heating energy consumption, 30–40 °C can be selected as the optimal reaction temperature in the actual reaction process. So the optimal conversion conditions of this process are as below: a molar ratio of Ca(OH)2/(NH4)2SO4 is 1:1.2, the Ca(OH)2 concentration is 37 g L−1, the reaction time is 30 s, a distance between the stator and rotor is 0.5 µm, the temperature is 30 °C, and the conversion efficiency of Ca(OH)2 is as high as 98.6%. This lime‐intensification sulfation process based on the RLFR has a Ca(OH)2 utilization rate as high as 99% and a reaction time as low as 30 s, providing the raw material basis for the regeneration of (NH4)2SO4 into NH3·H2O, which will be used for the desulfation of the next batch of spent lead paste after decomposition.In addition, the smaller distance between the stator rotors is beneficial to improve the sulfation rate of Ca(OH)4 product through the lime sulfation intensification process by the RLFR device under optimal experimental conditions. The obtained CaSO4 has a finer rod‐like structure , revealing the successful conversion of Ca(OH)2 into CaSO4. The residue Ca(OH)2 content has been titrated by the ethylenediaminetetraacetic acid (EDTA) titrationmethod through the dissolving of CaSO4 byproducts in excessive acetic acid  solution, the results indicate the average Ca(OH)2 content is around 2%, revealing the conversion rates of Ca(OH)2 are about 98%.Figure e Figure . In addi4 Figure . Two par4 Figure . The PXR2.32 as raw material and operated ten batches of continuous intensification desulfation and lime sulfation processes under optimal conditions ns Figure5a,b. Fi2.4To better evaluate the industrial implementation of the dual RFLR intensification desulfation process, we use 1 ton of spent lead paste as a unified starting point to calculate and compare the economic benefits of the traditional desulfation process and the new intensification desulfation process, as shown in Figure 3In summary, we propose an ultra‐fast recyclable and value‐added desulfation method for spent lead paste via dual intensification processes in which two newly introduced RLFRs for the enhancement of the desulfation process of spent lead paste and sulfation process of lime and mother liquor. The systematic investigation of the dual intensification processes parameters and the simulation calculation in a mathematical model of the experimental results demonstrate that the intensification desulfation process of lead paste and sulfation process of lime are controlled by internal diffusion. RLFR device greatly promotes the improvement of the mass transfer process and reaction rate through high‐speed grinding, pulverizing, and enhanced liquid dispersion processes, which achieve the desulfation rate and conversion rate of the new process are 99.7% and 98.6% in 10 and 30 s, far higher than the 99.4% and 91.7% of the traditional process in 40 and 30 min, respectively. Moreover, ten batches of continuous scale‐up experiments also fully confirm the stability of the process. The new RLFR intensification process has brought huge technical and economic benefits in space‐time yield and raw material conversion rate, which can significantly reduce the investment cost and improve net profits of the recycling technologies, promoting the development of the lead industry and recycling of the other related solid waste resources.4The experimental devices of desulfation of lead paste and sulfation of lime are shown in Figure 4)2CO3 solution was performed in an RLFR. The effects of different molar ratios of PbSO4 to (NH4)2CO3 in spent lead paste, reaction temperature, reaction time, (NH4)2CO3 concentration, and the gap between stator and rotor on desulfation rate were systematically studied.In the factor experiment, the desulfation process of 50 g spent lead paste (Chilwee Group) (Table 4)2SO4 solution (the previous step) obtained by desulfation of spent lead paste reacted with 28 g Ca(OH)2 in the RLFR to produce ammonia water and calcium sulfate. The effects of different molar ratios of (NH4)2SO4 to Ca(OH)2, reaction temperature, reaction time, lime concentration, and the gap between stator and rotor on lime conversion rate were investigated. Compared with the new process, the traditional desulfation process used a three‐necked reactor with a motor stirrer in the laboratory to optimize the process parameters in the desulfation of PbSO4 and sulfation of Ca(OH)2.(NH4)2SO4 reacted with Ca(OH)2 to generate ammonia solution and CaSO4. Then, the regenerated ammonia solution absorbed carbon dioxide that originated from the decomposition of desulfated lead paste to form (NH4)2CO3, which was recycled for the subsequent batch desulfation reaction. If the regeneration cannot be completed, the ammonia can be recovered from waste liquid by the simplified ammonia stripping technology. The regeneration ratio of (NH4)2CO3 is calculated by the following equationN represents the regeneration ratio of (NH4)2CO3; V1 represents the volume of the regenerated ammonia absorbing carbon dioxide to produce ammonium carbonate; C1 (mol L−1) represents the concentration of (NH4)2CO3; V2 is the volume of (NH4)2SO4 produced by desulfation of spent lead paste; C2 (mol L−1) is the concentration of (NH4)2SO4.After the desulfation and sulfation process, the obtained (NHThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


The strong interfacial interaction renders the material of good structural durability and enhanced reaction kinetics. Meanwhile, the enlarged and few‐layered MXene nanosheets can be easily obtained according to this interaction, providing a conductive network for sufficient electrolyte penetration and rapid charge transfer. As predicted, the VS2/MXene nanostructures exhibit an extremely low sodium diffusion barrier confirmed by DFT calculations and small charge transfer impedance evidenced by electrochemical impedance spectroscopy (EIS) analysis. Therefore, the SIBs based on the VS2/MXene electrode present first‐class electrochemical performance with the ultrahigh average initial columbic efficiency of 95.08% and excellent sodium‐ion storage capacity of 424.6 mAh g−1 even at 10 A g−1. It also shows an outstanding sodium‐ion storage capacity of 514.2 mAh g−1 at 1 A g−1 with a capacity retention of nearly 100% within 500 times high‐rate cycling. Such impressive performance demonstrates the successful synthesis strategy and the great potential of interfacial interactions for high‐performance energy storage devices.Constructing nanostructures with high structural stability and ultrafast electrochemical reaction kinetics as anodes for sodium‐ion batteries (SIBs) is a big challenge. Herein, the robust 2D VS 2/ Ti3C2Tx MXene nanostructures with the strong Ti─S covalent bond synthesized by a one‐pot self‐assembly approach are developed. The strong interfacial interaction renders the material of good structural durability and enhanced reaction kinetics, leading to the first‐class electrochemical performance of sodium‐ion batteries based on the VS2/MXene electrode.In this study, the robust 2D VS However, 2D materials are very easy to aggregate and self‐stack, which inevitably limits the sodium storage performance. Therefore, in the past decade, great efforts have been devoted to designing new 2D nanostructured electrodes for high‐performance SIBs.Two‐dimensional (2D) nanostructures have been extensively adopted as anode materials for sodium‐ion batteries (SIBs) due to their unusual physical/chemical properties.4, 5 Although it shows high electronic conductivity, tunable interlayer spacing, and abundant surface groups, its application in SIBs still suffers from the tendency to stack and the low intrinsic capacity. Moreover, it has been demonstrated that the ‐F and ‐OH terminal surface groups on MXenes could decrease the ionic conductivity, thus severely impeding the ion transport and restricting sodium storage. Many strategies have been proposed to address the above‐mentioned challenges, such as synthesizing the single‐/few‐layer MXenes, expanding the interlayer spacing, and constructing the MXene‐based composite materials.Recently, transition metal carbides, nitrides, and carbonitrides (MXenes), as emerging star 2D materials have gained much attention from researchers.12 Alt synthesized the composite of MXene nanosheets and black phosphorus quantum dots as the anode electrode for SIBs. The formation of P─O─Ti interfacial bonds accelerates the interfacial electron transfer, enabling fast and stable energy storage. Guo et al. reported the phosphorene/MXene nanoarchitecture and confirmed that the sodium diffusion kinetics in the phosphorene/Ti3C2F2 have been significantly enhanced. In particular, loading transition metal sulfides including layered sulfide materials (LTMD) and nonlayered materials on MXenes is of great significance to enhance the performance of SIBs. Because of the ability to undergo conversion or/and alloying electrochemical mechanisms, the sulfide/MXene composites almost all show an improvement in specific capacities and stabilities.Among them, MXene‐based composite materials could not only prevent the agglomeration/accumulation of MXenes, but combine the advantages of single components while making up for the relevant shortcomings to achieve improved performances for SIBs. For instance, Meng et al.33 synNevertheless, there are still a few important problems that need to be solved and understood. 1) The preparation for high‐quality few‐layered MXenes, especially in composites, is still a big challenge, which leads to insufficient exposure of the active site, resulting in the low concentration loading and poor utilization of active materials. 2) The weak interaction force between MXene and the second phase induces large electrochemical polarizations and slow interfacial charge transfer, limiting the high‐rate performance of SIBs. 3) The formation of an unstable solid electrolyte interface inevitably causes low initial columbic efficiencies (ICE) and rapid capacity degradation. 4) Among many studies of MXene‐based composites, few attempts have been made to understand the underlying mechanisms for sodium‐ion storage. Overall, there is still a long way to go for the widespread application of MXene‐based anodes for SIBs VS2 monolayer has the typical metallic feature and faster atom diffusion rate than graphite and MoS2. However, the limited specific surface area and volume variations of VS2 during charge/discharge processes lead to poor cycling stability and low reversible capacity of batteries. Lately, theoretical investigations have found that the integration of VS2 with MXenes could greatly enhance the electronic conductivity and adsorption strength, which is proposed to be a preferable anode material for SIBs. Therefore, it is very exciting yet highly challenging to fabricate the unique VS2/MXene heterostructure, then significantly accelerate the kinetics and achieve high‐performance SIBs. More recently, VS2/MXene hybrid was successfully produced by a conventional hydrothermal synthesis approach and applied in supercapacitors and Li‐S batteries. It is noticeable that the peculiar structure has a great impact on the electrochemical energy storage performance. Inspired by these results, considering the high requirements of SIB electrode materials, it is necessary to effectively optimize the composition and structure of VS2/MXene composites to obtain superior electrochemical sodium‐ion storage.As a representative member of LTMDs,47 VS22/MXene nanostructures are elaborately and successfully synthesized through the self‐assembly strategy of well‐dispersed few‐layer MXene nanosheets on VS2 without any surfactants followed by a freeze‐drying process and used as the anodes for SIBs for the first time. The enlarged, few‐layer MXenes can play the role of a 3D connected, conductive network and buffer matrix for VS2 to fully deliver its high capacity advantage. More importantly, it is found that there is the formation of strong Ti─S covalent bonds between two components, which is favorable to the efficient interfacial electron transfer and improves the surface capacitive contribution for sodium storage. In addition, according to the DFT calculations, the sodium absorption energy and diffusion energy barrier for the interface have been significantly decreased, enabling ultrafast kinetics for electrochemical reactions. Consequently, the 2D VS2/MXene nanostructures electrode exhibits an ultrahigh ICE (95.08%), impressive rate capacity of 424.6 mAh g−1 at 10 A g−1, and outstanding cycling stability of 514.2 mAh g−1 with capacity retention nearly 100% within 500 cycles at 1 A g−1. These results highlight the role of 2D nanostructured electrodes for high‐performance rechargeable batteries and demonstrate the importance of interface interaction in effectively boosting the kinetics and stability of electrodes.Hence, in this work, the 2D VS22/MXene nanostructure is described in Figure 2 and MXene, followed by the self‐assembly of MXene on VS2 nanosheets. On one hand, the VS2 nanosheets are achieved through the hydrothermal method with later ultrasonic treatment in DI water. On the other hand, MXene nanosheets are synthesized by selectively etching the pristine Ti3AlC2 in LiF‐HCl concentrated solution and subsequent manual‐shaking approach. Finally, the VS2/MXene nanostructures are fabricated by assembling MXene onto VS2 nanosheets (weight ratio is 1:5) driven by Van Der Waals forces with the assistance of longtime stirring. Zeta potential properties of nanosheets are first measured and shown in Figure 2 and MXene are negatively charged in the pH range of 4–10. Interestingly, at a pH of 7, the values of zeta potential of VS2 and MXene are almost the same −18.37 and −18.10 mV, respectively. Noted that the pH for the solution used in the synthesis for hybrid materials is very close to 7, which is more favorable to the interaction between two components. Figure 2 and MXene. The strong Tyndall scattering effect can be clearly observed when directing a side laser light on the colloids, demonstrating the good and stable dispersion of materials in aqueous media. Noteworthy, the mixture of two colloids of VS2 and MXene shows a homogeneous dispersion with no apparent precipitation due to the electrostatic repulsion, indicating the good compatibility between the two types of nanosheets. Moreover, the zeta potential of VS2/MXene is measured as −20.3 mV, which means that the hybridization does not change the overall state of the surface charge.The overall fabrication process of the VSFigure 2 microflowers are formed from a large number of nanosheets with a thickness of ≈20–50 nm and a diameter of ≈1.5 µm. Besides, the SEM image for MXene  and (104) plane of MXene, respectively. The SAED pattern , and a weak peak below 10° is attributed to the MXene, which may be due to either the small amount or the low crystallinity of MXene in the composite. Figure 2 and MXene signals are well identified in the hybrid. Meanwhile, in comparison with pristine MXene, the ratio for the D band and G band of VS2/MXene decreases from 0.80 to 0.67, suggesting a higher graphitization degree. Besides, the XPS spectrum is utilized to visualize the surface chemical compositions of materials. Figure 2/MXene nanostructures, where the characteristic peaks of V, S, Ti, C, O, and F can be clearly seen in the chart. The high‐resolution V 2p spectrum of VS2/MXene , Ti─S bond and S2+, respectively. Compared with the S 2p spectra of pristine VS2  and slightly lower than that of pure MXene (57.2 m2 g−1). The slight decrease relative to pure MXene is attributed to the successful loading of VS2 and the large increase compared with pure VS2 demonstrates the great advantages of few‐layer MXenes. Furthermore, the pore size distribution curves  and 8 times of individual MXenes (53.6 mAh g−1). This result suggests the excellent rate performance for VS2/MXene nanostructures, which could be further verified by a capacity of 628 mAh g−1 as the current density goes back to 0.1 A g−1. Besides, the shape of corresponding galvanostatic charge–discharge curves  curves are first investigated with a voltage window of 0.1–3.0 V at 0.1 mV s−1, and the results are shown in Figure  Another two peaks at 1.57 and 1.05 V are ascribed to the multi‐step sodium‐ion insertion of VS2 and MXene, which then shift to the lower potential of 1.56 and 0.82 V in the following cycles, respectively. During the anodic scan, peaks between 1.3 and 2.3 V can be observed, corresponding to the sodium‐ion extraction out of the materials. The well‐overlapped curves for 2–5 cycles demonstrate the satisfied electrochemical reversibility of the VS2/MXene nanostructures during charge/discharge processes. Furthermore, the CV curves at different scan rates from 0.1 to 5.0 mV s−1 for VS2/MXene are recorded and depicted in Figure 2/MXene in the dimethylether (DME) electrolyte. Besides, a related analysis can be conducted according to the relationship between the measured peak current (i) and scan rate (v) described by the following equationa and b are both adjustable values. Therefore, the slope of the corresponding log(v)‐log(i) plots 2 stacked in the vertical direction through the interaction of Van De Waals forces. The monolayer VS2, Ti3C2F2/Ti3C2(OH)2, and their heterostructures are all investigated. Figure 2/Ti3C2F2 and VS2/Ti3C2(OH)2 present a very high absorption energy toward sodium atoms, which is much better than that of the individual components. This confirms that the heterostructure is beneficial for sodium storage. Meanwhile, the energy barrier for sodium diffusion is also investigated and shown in Figure 2/Ti3C2F2 interface is lower than that of diffusing along the surface of monolayer VS2 and Ti3C2F2. The difference between the interface and Ti3C3F2 is 0.157 eV (DFT accuracy is around 0.1 eV), so it can be considered that the formation of the interface enhances the sodium diffusion kinetics. Moreover, the energy barrier of sodium diffusion in VS2/Ti3C2F2 (0.0464 eV) is much lower than the previously reported sodium diffusion barrier in other 2D heterostructures, such as BP/Ti3C2F2 (0.1729 eV) and MoSe2/MXene (0.095 eV). This low energy barrier ensures fast ions movement during the sodiation/desodiation process, leading to the high rate performance for SIBs. Besides, the energy barrier for Ti3C2(OH)2 is calculated as 0.0216 eV, whereas the barrier for the interface of VS2/Ti3C2(OH)2 is 0.0823 eV. Due to the accuracy of the DFT, the difference of only 0.06 eV can be considered that there is little difference in the diffusion of sodium atoms between the two structures and the sodium spreads out very quickly in both. Furthermore, the galvanostatic intermittent titration technique (GITT) profiles shown in Figure 2/MXene during the whole sodium storage reactions compared with the individual components, further verifying the fast diffusion of sodium. Next, the electron excursion is determined by the calculation of Bader charges. The side view for VS2/Ti3C2F2 structure is shown in Figure 2 and Ti3C2F2, the interface shows the electron‐rich state of V and Ti, as well as the electron‐lack state of S, C, and F. Consequently, the charges for S ions, which could transfer to the Ti ions through Ti─S bond, further confirming the strong interaction between VS2 and MXene. Meanwhile, the charges for C atoms could transfer to V ions through F ions. Therefore, the interaction between VS2 and Ti3C2F2 can be depicted by the following reaction .2/Ti3C2(OH)2 structure is also performed and the results are shown in Figure 2/Ti3C2F2 and the interaction between VS2 and Ti3C2(OH)2 can be given as follows.The DFT calculations are further performed to investigate the absorption and diffusion characteristics of sodium in the VSs forces.71 The32/MXene nanostructure is synthesized through the simple self‐assembly strategy to prevent the agglomeration/accumulation of MXenes and protect VS2 to fully deliver its high capacity advantage. The enlarged, few‐layer MXene nanosheets are tightly anchored into the VS2 nanosheets with the strong interface interaction as well as the formation of Ti─S covalent bond, endowing the material with high structural durability and accelerating the electrochemical activity. Besides, the VS2/MXene anode for SIBs possesses a low sodium diffusion barrier (0.0464 eV) calculated by DFT, small charge transfer impedance, and high surface capacitive‐dominated contribution, demonstrating the ultrafast kinetics of electrode during the charge‐discharge processes. As a result, the fabricated SIBs exhibit ultrahigh average ICE, excellent rate capability, and prolonged cycling stability. The successful synthesis strategy would enable the VS2 for prospective anode materials for SIBs and can be extended to other novel nanostructures for impressive electrochemical performance.In summary, a unique 2D VSThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


The family owes its potential to its ferromagnetically (FM) ordered MnBi2Te4 septuple layers (SLs). However, the QAHE realization is complicated in MnBi2Te4 and MnBi4Te7 due to the substantial antiferromagnetic (AFM) coupling between the SLs. An FM state, advantageous for the QAHE, can be stabilized by interlacing the SLs with an increasing number n of Bi2Te3 quintuple layers (QLs). However, the mechanisms driving the FM state and the number of necessary QLs are not understood, and the surface magnetism remains obscure. Here, robust FM properties in MnBi6Te10 (n = 2) with Tc ≈ 12 K are demonstrated and their origin is established in the Mn/Bi intermixing phenomenon by a combined experimental and theoretical study. The measurements reveal a magnetically intact surface with a large magnetic moment, and with FM properties similar to the bulk. This investigation thus consolidates the MnBi6Te10 system as perspective for the QAHE at elevated temperatures.The recent realizations of the quantum anomalous Hall effect (QAHE) in MnBi 6Te10 is a promising route toward topological qubits, ultra‐low‐power electronics and spintronics. The most fundamental ingredient is ferromagnetic order; however, in pristine MnBi6Te10, the consecutive Mn layers interact antiferromagnetically. Here, it is demonstrated how Mn and Bi intermixing between their respective atomic positions may induce robust ferromagnetism.Stabilizing the quantum anomalous Hall effect in a magnetic topological insulator such as MnBi The vision of observing Majorana fermions and implementing topological qubits at superconductor/QAH insulator interfaces, ultra low‐power electronics, and applications in spintronics has ignited substantial experimental efforts in this direction. Yet, hitherto the QAH effect (QAHE) has only been demonstrated in the (sub‐)Kelvin range. The experimental realization of the QAHE is complicated by several simultaneous requirements to a candidate system: The Dirac point (DP) of the parent TI should be well within its bulk band gap; the chemical potential has to be tuned to the DP; the introduced magnetic subsystem should lead to a substantial surface ferromagnetism to open a large exchange gap at the DP; and the material's bulk should remain insulating.Theory provides a seemingly straightforward avenue toward novel quantum effects such as the quantum anomalous Hall (QAH) effect,4, 5xx2 − Te3, which consist of van‐der‐Waals coupled quintuple layers 2Te4)(Bi2Te3)n , whose functional constituents are (MnBi2Te4) septuple layers (SLs) with the central sheet of FM‐ordered Mn atoms, separated by n Bi2Te3 QLs, offer several advantages. Whereas for QL termination, angle‐resolved photoemission spectroscopy (ARPES) measurements on MBT1 and MBT2 yield similar results to 2Te3, without a discernible DP, there is a DP within the bulk gap for the SL termination. Also, the topmost sheet of the ferromagnetically arranged Mn moments should strongly couple with the topological surface states (TSS), albeit a thorough spectroscopic investigation of the surface magnetism is still pending. As a result, a much higher TQAH = 1.4 K is achieved in MnBi2Te4. This is despite the fact that MnBi2Te4 is suboptimal due to its antiferromagnetic (AFM) order (TN = 24 K) and a complex layer‐number dependence of the quantization effects, with an odd number of SLs required to realize the QAHE.The intrinsic magnetic topological insulators , an intrinsic magnetic TI and QAHE candidate.Yet, the potential of the other MBT6Te10 crystals by using bulk‐sensitive superconducting quantum interference device (SQUID) magnetometry and surface‐sensitive X‐ray magnetic circular dichroism (XMCD). The clear FM characteristics seemingly contradict the weak AFM coupling anticipated by our density functional theory (DFT) calculations for the atomically ordered compound. This disagreement is resolved by including the experimentally determined Mn substoichiometry and Mn/Bi site intermixing into account. Our calculations pinpoint that the magnetic coupling can be tuned toward ferromagnetism by appropriate intermixing already in MnBi4Te7 and even MnBi2Te4. Considering the intermixing patterns in our MnBi6Te10 samples and those reported showing no ferromagnetism, we rationalize their differing magnetic behavior. Our results demonstrate that carefully engineered intermixing can accomplish a robust FM order and, therefore, is the key toward enhanced QAHE properties in the MBTn family of intrinsic magnetic topological insulators.We confirm the FM state both in the bulk and on the surface of MnBi22.16Te10 crystals were grown by slow crystallization from a melt . Again, this echoes our earlier published single‐crystal structure refinement of Mn0.73(4)Bi6.18(2)Te10 by X‐ray diffraction, where we systematically showed that Mn‐substoichiometry is determined by the Mn/Bi intermixing. Both features are also present in Mn0.85Bi2.10Te4 and Mn0.75Bi4.17Te7. To facilitate perception, we denote our samples as MnBi6Te10 in the following text, keeping in mind that they are in fact substoichiometric.A series of energy‐dispersive X‐ray spectroscopy (EDX) point measurements on individual crystals extracted from the ingot demonstrated a compositional range between Mn: 5.0, Bi: 36.6, Te: 58.4, and Mn: 4.2, Bi: 37.1, Te: 58.7 (in at%). Our samples were thus consistently more Mn‐deficient than expected from the nominal chemical formula MnBifraction,31 whe6Te10 crystals  see Figure 6Te10 with a sequence of one SL and two QLs  which crystallized in the rhombohedral space group R3¯m (No. 166) with the unit cell lattice parameters a = 4.36778(8) Å and c = 101.8326(6) Å. To stabilize a further Rietveld refinement, the EDX compositions  were introduced as constraints , the acquired powder diffraction data did not allow us to settle in for just one particular intermixing model with a statistically unequivocal quantification. The refined Mn content is also strongly dependent on whether cation vacancies are allowed in the refinement. We opted for a structural solution without voids in the 3a and 6c positions. Despite the outlined uncertainties, all tested models with various composition constraints have in common that: 1) the Mn:Bi ratio in the 3a position in the center of an SL is close to 56:44; 2) the outer cation site of an SL (6c) contains up to 2% Mn; 3) the QL always accommodates some Mn (2–7% Mn) in the 6c cation sites. The presence of Mn in all cation positions accords with our earlier reported refinement on Mn0.81Bi6.13Te10 single crystals and is in contrast to the findings of Klimovskikh et al. Such subtle variations in intermixing patterns can dramatically impact the magnetic properties, as witnessed in the next subsection.This approach is exemplified on sample #2   in Figure 0Hc ≈ 32 mT at T = 2 K, and a finite remanent moment of mSQtot=(3.9±0.2)μB/Mn at zero magnetic field. The moment at 0.15 T is mSQtot=(4.2±0.2)μB/Mn.The magnetization curves n family  has a clear A‐type AFM structure, whereas MnBi4Te7 (n = 1) exhibits a more complex behavior, in which robust low‐temperature metamagnetic properties are established, which were shown to result from the competition between the uniaxial anisotropy K and the still sizable interlayer AFM interaction J. Finally, in MnBi6Te10 (n = 2), as well as in MnBi8Te13 (n = 3), the FM properties clearly dominate, with FM order at the significant temperatures of Tc = 12 and 10 K, respectively. Consistent with this observation, the spin‐flop transition found for MnBi2Te4 and MnBi4Te7 at fields of 3.5 T and 0.1–0.3 T, respectively, is absent in MnBi6Te10, and a magnetic moment of more than 4μB is observed already above 80 mT in the latter after a ZFC procedure.It is furthermore interesting to compare our results to analogously synthesized samples of the MBTy Figure . We obse2.3U for MnBi6Te10 neglecting the intermixing. For the interaction parameters, we have used the Slater integrals in Table We have first performed fully relativistic DFT calculations based on the Generalized Gradient Approximation (GGA)+6Te10, with its lattice parameter c > 100 Å, would require a prohibitively long computational time. Instead, here we aim to learn the effects of Mn/Bi intermixing on magnetism via the simpler models of MnBi2Te4 and MnBi4Te7. The latter case is more representative of MnBi6Te10, since it contains both SLs and QLs, and is discussed in detail below, while the former is presented in Section 2Te4: even though its defect‐free form has the strongest AFM coupling between Mn in the consecutive SLs, based on our calculations, intermixing can induce the FM order between the SLs even in this compound  of both the SLs and the QLs. MnBi4Te7 provides the minimal framework to explore whether this enlarged configuration space can yield variations in the experimentally observed ground states. We have performed scalar‐relativistic calculations for various intermixing patterns and magnetic orders in a 2 × 1 × 2 supercell of MnBi4Te7 , we construct models (S1 to S15) with different Mn/Bi antisite defects, all of them globally stoichiometric and having a 50% fraction of Bi atoms in the 3a Wyckoff site. Notice that this concentration is close to the outcome of our Rietveld refinements (≈44%). The models differ in the positions occupied by the antisite Mn atoms and can be classified into three categories. In the first category, the antisite Mn atoms reside in the 6c position of only the QLs (S1 to S4). Similarly, in the second category antisite Mn occupy the 6c positions of only the SLs (S5 to S8). In the third category, the Mn atoms are distributed over the 6c positions of both the QLs and the SLs (S9 to S15).As 7 Figure 3. In ad1 and AFM2 models, the Mn spins in the two consequent 3a positions are oppositely coupled. The coupling between the 3a site and all intermixed Mn neighbors in the 6c site(s) is either AFM (AFM1) or FM (AFM2), respectively. The AFM3 model realizes parallel spin arrangement in the 3a sites, while they couple AFM with the Mn defects in all 6c positions.For each structure model, we consider four possible magnetic arrangements: a fully spin‐polarized FM order and three different AFM models sketched in Figure 4Te7 as compared to the fully spin‐polarized FM state. For a given model, if any AFM model obeys EAFM − EFM < 0, we conclude that antiferromagnetism is preferred. On the other hand, if all AFM models fulfill EAFM − EFM > 0, we define the fully FM phase as the ground state.Figure either the SLs or the QLs, show an AFM configuration as the lower energy state. This preference reverts markedly if the antisite Mn distributes over the 6c positions of both the QLs and the SLs: five out of the seven constructed structure configurations of this category prefer the FM phase. A closer look reveals that the preference for the FM state is particularly prominent in those structure models (S12 to S15), in which the Mn cations occupy the nearby 6c positions not separated by the 3a positions—namely, when continuous magnetic exchange pathways exist between the antisite Mn ions.A clear trend in the magnetic order as a function of the underlying Mn/Bi intermixing pattern can be established. All but one models of the first two categories, where the intermixed Mn cations occupy c positions of the QLs—a situation experimentally found in ref. [c of the QLs and SLs, which is the case of our samples according to our structure refinements, our calculations identify the FM phase as the ground state.These results establish a strong correlation between the magnetic structure and the Mn distribution along the stacking direction. When an antisite Mn is located only in one of the two 6 in ref. —our calc2.4N4, 5 edges  measurements in the total electron yield (TEY) mode, which is element specific and has a probing depth on the nanometer scale. Measurements at the Bi L2, 3 edge. Due to the shallow escape depth, the topmost SL contributes the most to the signal. However, even for SL termination, the FM Mn sheet is buried about 0.55 nm below the surface, and significantly more for QL termination. Therefore, probing depth effects have to be considered, when interpreting the ordered magnetic moments obtained with XMCD . Within the precision allowed by the T increments of 2 K, the transition temperatures at surface and bulk are consistent. We remark that the transition behavior as observed with SQUID and XMCD could differ somewhat due to the different measurement protocols: For XMCD, each point in Figure 0H = 3 T and back to remanence. In contrast, in SQUID measurements a conventional FC protocol at 10 mT was used.Finally, in Figure 2.5 , rather than relying on oversimplified approximations, we adjust most of the MLFT parameters to the data and obtain 10Dq = 0.06, 10DqL = 2Tpp = 1.9, Δ = 1.1, Udd = 4.0, Upd = 5.0, Veg=1.3 and Vt2g=0.65 . For the SO coupling constants, we use the Hartree‐Fock values, whereas the Slater integrals are calculated based on DFT in the local density approximation  dominates with 71%, there is significant charge transfer from the Te ligands, resulting in a 27% contribution of d6L_ to the ground state μB/Mn, which is compatible with sample #4 within the error.The MLFT analysis yields the s Figure . FinallymXMspin is to analyze the XMCD L3 peak asymmetry, which avoids the problems arising from uncertainty in p due to the overlap of the L3 and L2 peaks (Section mXMspin=(2.55±0.25)μB/Mn, which is about 10% larger than the sum rule result.An alternative way to obtain 5 configuration  positions of the SLs might contain slightly canted Mn, as well as a few percent of paramagnetic and possibly even AFM (with respect to the 3a positions) Mn, see Section a positions) is buried about 0.55 nm below the surface. At the same time, there is Mn in 6c positions closer to the surface, both in the very same SL and in QLs, which can terminate the surface in different parts of the sample. Therefore, already for a mean probing depth (MPD) of about 1 nm, the contribution of the FM Mn would be notably suppressed. MPD values between 1 and 2.5 nm have been reported for this photon energy range. Given the presence of the heavy elements Te and Bi, which might even further attenuate the escaping electrons, an MPD close to 1 nm is not unrealistic and the involvement of probing depth effects is well conceivable.It is important to keep in mind that the indicated errors of the XMCD results only take into account statistical fitting and background estimate effects. However, the shallow probing depth can further bias the outcome (Sb2Te3)n (MSTn) becomes relevant. The FM order is more dominant even for n = 0 in MSTn and MBSTn, and, as widely accepted by now, is driven by the Mn/Sb intermixing. This phenomenon is much stronger in MSTn than in MBTn since it is facilitated by closer atomic radii of Mn and Sb. On the other hand, the impact of the intermixing‐induced FM state in MnSb2Te4 on its band topology is still under ongoing debate. Intrinsic p‐type doping in MST0 hampers clear‐cut spectroscopic observations of the possible surface states and, thus, an ultimate conclusion about its topological nature. Furthermore, QAHE realizations in the MSTn have not been reported. In fact, ref. [n materials, but focuses on the MST0 instead, because the necessary intermixings were not accessible by the bismuth analog at that time. So far intermixing in the MBTn has been discussed mostly in the terms of its influence on the Dirac‐point gap, while the consequences for the magnetism are quite unclear. En route to understanding the broader role of intermixing, a recent study reveals its crucial influence on the magnetic coupling in MnBi2Te4, and our current work pinpoints the particular antisite defects that enhance (or suppress) the local FM coupling in the MBTn series.In this context, a comparison between the MBTct, ref.  argues t6Te10 is often found by X‐ray spectroscopy, the related intermixing has been scrutinized only in ref. [a site. On the other hand, there are also substantial differences: The mixed 3a occupancy is more pronounced in our sample, in which we find 56% Mn (and 44% Bi), than in the sample studied in ref. [c positions is distinctly different: We observe a higher Mn concentration in both 6c sites of the QLs, that is, up to Bi1.86Mn0.14Te3 versus Bi1.92Mn0.08Te3 in ref. [a site in our crystals is more Mn‐depleted, so that these “stray” Mn atoms, which find no space on the 3a site, disperse over the entire layered stack by occupying 6c sites. In accordance with our theoretical deliberations in Section 1 to S4), the less pronounced intermixing and the presence of swapped Mn only in one of the two 6c sites of the QLs in the samples of ref. [We have established that the FM properties of our crystals are conditioned by the underlying cation intermixing. It is instructive to examine whether this relationship holds true for the other published works. Whereas Mn deficiency in MnBitroscopy,25 the in ref. . On the  in ref. , which h in ref. , and up  in ref. . In gene of ref.  accords 2Te4 and MnBi4Te7. It has been argued that such defects provide an effective way to release a lattice strain effect which occurs within the septuple layer of MnBi2Te4 due to a mismatch between the MnTe and Bi2Te3 structure fragments. Second, as argued below, variations in the intermixing patterns of MnBi6Te10 samples produced by different groups may stem from subtle differences in the synthetic procedures, pointing to the relevance of finite temperature effects for the relative stability of different defects. This would be of no surprise since these compounds are formed at elevated temperatures and are metastable at room temperature as we have previously shown.The question of why intermixing takes place and which kind of defects are more likely to occur is evidently very relevant and, at the same time, a complex one. First, recent literature has shown that antisite cationic defects have the lowest formation energy and are energetically favorable to form in both MnBiMnBi4Te7.73 It 6Te10 to those of refs. [2 crystals  may be a blessing when it comes to optimizing an MBTn material's system for the QAHE device fabrication.Comparing our growth conditions . The ampule was subsequently cooled down to 858 K at a rate of 1 K h−1, tempered for 14 days and then quenched in water. Platelet‐like MnBi6Te10 crystals  were mechanically separated from the obtained ingot.Pre‐synthesized, phase‐pure MnTe and Bi1 radiation (λ = 154.056 pm). The phase composition of the polycrystalline ingot and individual crystals was estimated by Le Bail or Rietveld methods in JANA2006. The preferred orientation of the crystallites was described by March–Dollase corrections, the roughness for the Bragg–Brentano geometry was accounted for by the Suorti method.Powder X‐ray diffraction data were collected on an X'Pert Pro diffractometer  with Bragg–Brentano geometry (featuring variable divergence slits) operating with a curved Ge(111) monochromator and Cu‐KαJANA2006.78 TheUa = 2 − 5 kV. The composition of selected single crystals was determined by semi‐quantitative energy dispersive X‐ray analysis at 20 kV acceleration voltage.Scanning electron microscopy (SEM) was performed using a SU8020 (Hitachi) equipped with a X‐MaxN (Oxford) Silicon Drift Detector (SDD) at M per Mn atom, a precise knowledge of the sample mass is important. Samples #1 and #4 have ≈10 times smaller mass than sample #2, increasing the error of M. Nevertheless, the data for all four samples agreed well with each other , which was the quantity obtained from the SQUID magnetometry measurements, is referred to as mSQtot.Field and temperature dependent magnetization studies were performed using a quantum design SQUID magnetometer equipped with a vibrating sample magnetometer (VSM) option (MPMS3). The magnetization data on samples #1–#4 were normalized to the real compositions determined via EDX. To obtain the absolute magnetization CW and the temperature independent susceptibility χ0 for the low‐mass samples MnBi6Te10.A setup made of two half‐cylindrical quartz rods fixed with a small quantity of GE varnish to the main quartz VSM sample holder was designed to ensure an alignment of the crystals such that the external magnetic field was applied perpendicular to the crystal surface. Note that this setup, however, resulted in a rather temperature‐independent (at not too low temperature) but non‐negligible background contribution to the magnetic susceptibility, hindering a reliable extraction of the Curie–Weiss constant θU were performed with the parametrization of Perdew, Burke, and Ernzerf, using the full localized limit for the double‐counting correction with U = Udd (the latter as obtained in Section J=(Fdd(2)+Fdd(4))/14, with Fdd(2) and Fdd(4) the Slater integrals for the initial states presented in Table 6Te10, a mesh of the Brillouin zone having 14 × 14 × 14 subdivisions was used. The magnetic anisotropy energy was calculated in the AFM state based on a mesh having 10 × 10 × 10 subdivisions. The Mn 3d occupancy and the spin projection presented in the main text correspond to the gross projections. For the calculations of MnBi4Te7 based on a 2 × 1 × 2 supercell, a mesh with 6 × 12 × 2 subdivisions was used.Fully relativistic DFT calculations based on the GGA +  Ernzerf,38 usi and at the high‐field diffractometer at the UE46 PGM‐1 beamline, BESSY II. The single crystals were glued with conducting silver epoxy onto Cu sample plates and mounted on the cold finger of a helium flow cryostat. Prior to the measurements, each sample was mechanically cleaved in the fast‐entry chamber at a pressure of ≈10−9 mbar to expose a pristine surface. The sample was then transferred into the spectroscopy chamber with a pressure in the 10−11–10−10 mbar range.The XAS and XMCD measurements were performed using the high‐field cryomagnet end station HECTOR of the BOREAS beamline at the ALBA synchrotron radiation facility79 andI0). Circularly polarized light was used at both beamlines. The area probed by the beam at both facilities (about 120×80μm2) was much smaller than the sample size.The measurements were carried out in the TEY or FY mode at magnetic fields of up to 6 T and at various temperatures in the 3.5–35 K range. The temperature was calibrated with a thermal sensor mounted at the sample position before the experiment. Especially below about 5 K, the actual sample temperature crucially depends on the thermal contact, increasing its error as compared to higher temperatures. The spectral intensity was normalized by the incoming photon intensity /2. To cancel out any experimental drifts, for each data set eight spectra were measured in a row by altering the X‐ray polarization. Finally, the magnetic moments measured with XMCD are marked with the subscript XM, for example, mXMtot.The raw XAS spectra were scaled with respect to each other to have the same intensity at energies far from the resonances to obtain  The Brillouin zone was sampled by a 2 × 2 × 2 k‐point mesh. The exchange‐correlation potential was treated in LDA, with the scalar relativistic functional according to ref. [R3¯m (166), a = 4.37 Å and c = 101.83 Å, slightly distorted octahedral Mn coordination with Mn−Te bond length of 3.00 Å  was used. Wannier orbitals were obtained as input for MLFT by downfolding to a basis set of Mn 3d, Te 5p, and Bi 6p orbitals in an energy window from −6 to 3 eV including an exponential decaying tail with a decay of 1 eV at the boundaries of the selected energy range.As a starting point to obtain input parameters for the MLFT modeling, self‐consistent DFT in the linear density approximation is sufficient, which was performed using the FPLO package.81 The to ref. . The exp to ref.  was used within the CI scheme considering the nominal 2p63d5 (Mn2 +) configuration and two further charge‐transfer states d6L_ and d7L_2. The spectral contributions from the split ground‐state terms were weighted by a Boltzmann factor for T = 2 K. The mean‐field effective potential was modeled by an exchange field estimated from the Tc of 12 K. Instrumental and lifetime effects were taken into account by a Gaussian broadening of 0.35 eV (FWHM) and an E‐dependent Lorentzian profile of 0.15 − 0.35 eV (FWHM).The MLFT calculations were performed using the Quanty package84 witFdd(2) and Fdd(4) were scaled up by 8% for the final state, improving the agreement to experiment. SO coupling constants were kept to the Hartree‐Fock values.Δ=E(dn+1L_)−E(dn), Udd and Upd were directly fitted to the experimental spectra, keeping Udd/Upd = 0.8. Experiments involving charge‐neutral excitations such as XAS are only weakly sensitive to Δ, Udd, and Upd. In this particular case XAS and XMCD spectra were fitted simultaneously, which substantially mitigated these kind of problems. The results were in good agreement with values reported for As and Mn‐doped Bi2Se3 and Bi2Te3. The other MLFT input parameters were estimated from DFT, and their values were subsequently adjusted to reproduce the experimental spectra. To simplify the calculation, instead of the trigonal Cv3 the work was done in Oh symmetry, with the C4 octahedral axes along the Mn−Te bonds, which had a negligible impact: The simplification here neglected the splitting of the t2g orbitals, which is tiny compared to 10Dq < 100 meV, which in turn is smaller than the experimental resolution.The Slater integrals for the MLFT calculations were obtained by DFT, where k values.47Δ=Ep and q are the XAS intensity differences (Ileft − Iright) integrated over the L3 edge and the entire L2, 3 region, respectively (Figure I, after background correction (Section L2, 3 to yield r. 〈Tz〉 is the expectation value of the intra‐atomic magnetic dipole operator which is −0.0002ℏ and hence negligible (Section nd the MLFT value of 5.31 (Section C is a correction factor, which takes into account the considerable overlap of the L3 and L2 contributions for light transition metals. A value of C = 1.4 (Section mXMspin was also obtained by a comparison of the experimental XMCD asymmetry at the L3 peak to the theoretical one calculated from MLFT spectra of comparable line width.The XMCD sum rules yield: Section  was usedThe authors declare no conflict of interest.L.F., E.K., and A.I. conducted and analyzed the experimental work related to crystal growth, XRD, and EDX. B.R., L.T.C., and A.U.B.W. performed and analyzed bulk magnetometry measurements. J.I.F. and J.v.d.B. performed and analyzed density‐functional calculations aimed to understand the effects of intermixing. A.T., V.B.Z., T.R.F.P., P.G., M.V., and E.W. performed the XAS measurements. A.T. and V.B.Z. analyzed the XAS data with input and support from T.R.F.P., P.K., S.H., H.B., R.J.G., M.W.H., and V.H. V.B.Z., A.I., and V.H. conceived and supervised the project with input and support from F.R., J.v.d.B., B.B., and A.U.B.W. All authors contributed to the discussion and writing of the manuscript.Supporting InformationClick here for additional data file.


This interface is shown not to contribute to carrier transport, but to contribute to carrier concentration filtering due to carrier injection or recombination. The Seebeck coefficient of the bi-phase material is shown to be dependent on the direction carriers are crossing the interface. The interface effect mainly results from a modification of charge carrier concentrations in the semiconductor.Thermoelectricity is a well-known effect that can be used to convert heat energy into electrical energy. However, the yield of this conversion is still low compared to current photovoltaic technology. It is limited by the intrinsic properties of materials, leading to intensive materials science investigations for the design of efficient thermoelectric (TE) materials. Interface engineering was shown to be a valuable solution for improving materials’ TE properties, supporting the development of multiphase TE materials. In particular, interfaces have been suggested to promote the increase of the Seebeck coefficient of materials without significantly impacting their electrical conductivity through the so-called energy filtering effect. This work aims at determining experimentally the effect of a metal/semiconductor interface exhibiting an ohmic character on the effective Seebeck coefficient of multiphase materials, focusing on the  However, their yield of converting waste heat into electrical power is still currently low (~ 10%) as it dependents on interrelated material properties exhibiting contrasting effects on TE properties8. Indeed, the conversion efficiency of a given TE material is determined by the figure of merit ZT = T × S2σ/κ 7. Higher the ZT, higher the TE efficiency, meaning that efficient TE materials should exhibit high Seebeck coefficient, high electrical conductivity, and low thermal conductivity. However, S and σ are in general inversely proportional, and the electronic component of thermal conductivity increases with electrical conductivity, limiting the TE efficiency of materials in general.Worldwide efforts for the development of efficient thermoelectric (TE) devices is motivated by their ability of safely produce electricity from waste heat without CO10 and PbTe12, and half-Heusler compounds13, for example, aiming at discovering the most effective TE material. Though, multiphase materials were shown these past 10 years to exhibit higher ZT than single-phase materials19, allowing to reach ZT > 2 in some cases22, suggesting the possibility of reaching the limit ZT = 3, considered as the condition for the production of efficient devices, matching application requirements23. Usually, TE properties of multiphase materials are considered to be higher than single-phase materials due to the combination of increased phonon scattering at interfaces, increased Seebeck coefficient due to interface-mediated energy filtering, and increased electrical conductivity due to modulation doping. The effect of nanostructuration and nano-size phase inclusions on thermal conductivity has been directly proved thanks to heat conduction measurements30, and interpreted as due to a modification of the phonon component of thermal conductivity32, allowing the ZT numerator to be reduced. The increase of the power factor PF = S2σ (ZT denominator) in multiphase materials is often attributed to the effect of interfaces, in particular in the case of S increase, generally interpreted as the energy filtering effect of the interface between the matrix and inclusions34 or grain boundaries39.Consequently, a wide material panel is being explored for TE power generation applications, including the compounds GeTe42 and numerical calculations44 showing a possible increase of S due to interface-mediated energy filtering, the effect of interfaces on S still needs further experimental investigations in order to improve TE property engineering of multiphase materials. For example, the dependence of S with interface area in multiphase materials exhibiting enhanced S has been poorly studied45, and some experimental results shown no modification of S with the presence of interfaces48. Furthermore, the increase of S in multiphase materials has been shown as possibly due to a significant gradient of charge mobility resulting from the temperature gradient49. The junction between two conducting solids of very different charge carrier mobilities is thus expected to promote an increase of S that is not related to energy filtering of charge carriers.Despite theoretical modelsSeff) of a binary multiphase material, and to provide an experimental proof of Seebeck coefficient modification due to the interface. The experiments were performed around room temperature with complementary-metal–oxide–semiconductor (CMOS) compatible materials to match the requirements of CMOS-integrated TE device for energy harvesting in portable microelectronic circuits51.The present study aims at probing the influence of a single interface on the effective Seebeck coefficient  between the two materials before contact, an energy barrier is built (for holes in the present example) close to the interface after contact. This barrier is expected to scatter carriers of low energy, reducing the number of free carriers and increasing their average energy, leading to the increase of Seff (energy filtering effect)42.TE materials are usually Seff. This is nevertheless this type of interface that the present work aims at investigating, since it has been poorly studied yet and can be considered as a reference to be compared with other metal/semiconductor interfaces.A second type of metal/semiconductor interface can be built, corresponding to an ohmic contact. No major barrier is present at the interface in this case Fig. c, and ch5Ge3 and a polycrystalline thin film of p-type Ge. Mn5Ge3 is known to be n-type55 and to form an ohmic contact with Ge if Ge is p-type or a Schottky contact if Ge is n-type57. An n-type metal/p-type semiconductor interface allows possible carrier injection through the interface to be more easily evidenced. Mn and Ge were deposited at room temperature by magnetron sputtering on glass substrate58, and the Mn5Ge3/Ge bilayer was grown by reactive diffusion thanks to ex situ annealing under vacuum (P ~ 10−7 mbar) at 400 °C for 10 min60. The use of thin films allows the study of a single interface, as well as the volume of the phases to be precisely controlled. The sputtering technique guaranty a control of the film thickness below 5 nm, as well as a low level of contamination, below 1019 cm−3 in our case61. Furthermore, structural and electrical characteristics of the Mn5Ge3/Ge interface is generally considered to be well reproducible through reactive diffusion fabrication63.The chosen interface is the contact between a polycrystalline thin film of the metallic compound MnSeff of multiphase materials can be probed through two types of geometry: (1) the parallel-phase geometry, and (2) the serial-phase geometry. In our case, the parallel-phase geometry corresponds to the Seebeck coefficient Seff = ΔV/ΔT measured on the bilayer Mn5Ge3/Ge according to Fig. 5Ge3 and holes in Ge are moving towards the cold side of the bilayer. As presented in Fig. Seff should only depend in this case on the Seebeck coefficients (SGe and S5/3) and the electrical conductivities (σGe and σ5/3) of each material Ge and Mn5Ge3 if the interface has no effect. However, due to carrier confinement or due to the formation of a third phase at the interface for example, the interface can act as an additional carrier channel. In this case, the signature of the interface effect can be detected in Seff, as the equation given in Fig. 5Ge3/Ge interface64. In the serial-phase geometry  in the temperature gradient (in the colder region or in the hotter region), exhibiting nonreciprocity versus the temperature gradient direction65.The influence of an interface on the eometry:  the paraeometry:  the paraeometry:  the paraeometry:  the paraeometry:  the paraeometry:  the paraeometry:  the paraeometry:  the paraeometry:  the paraeometry:  the paraT) on a 100 nm-thick Ge (σGe) film and on a 160 nm-thick Mn5Ge3 film (σ5/3) deposited on glass substrate. σGe increases with temperature, corresponding to a semiconductor behavior for Ge, while σ5/3 decreases when the temperature increases, corresponding to a metallic behavior for Mn5Ge3, as expected. The conductivity is found to be about two orders of magnitude larger in the Mn5Ge3 metallic film compared to the Ge semiconductor film. Figure 5Ge3/Ge contact. The measurements were performed according to the Van der Pauw method on a bilayer Mn5Ge3/Ge sample, after removing part of the Mn5Ge3 film by chemical etching, making a 160 nm-thick step between a Mn5Ge3 upper terrace and a Ge lower terrace  of the 100 nm-thick Ge film and S5/3 (purple open squares) of the 160 nm-thick Mn5Ge3 film versus temperature. SGe > 0 and increases with temperature, while S5/3 < 0 and decreases when temperature increases. The opposite sign of SGe and S5/3 confirms that the Mn5Ge3/Ge ohmic contact  of the 160 nm-thick Mn5Ge3/100 nm-thick Ge bilayer measured according to the parallel-phase geometry. Figure Seff is in between of those of SGe and S5/3, but it is negative in the investigated temperature range and is close to S5/3. Seff-th calculated using the equation given in Fig. σGe and σ5/3 , while the thickness of the Ge layer was varied between 80 and 1380 nm in six different Mn5Ge3/Ge bilayers. All the measurements are close to Seff-th, the maximum difference between Seff and Seff-th at T = 300 K being 22% for example. Furthermore, no obvious dependence of the Ge thickness on Seff could be determined, Seff being closer to Seff-th in the sample with th = 880 nm than in the sample with th = 240 nm for example.5/3 Fig. a and of σGe Fig. a. Figure5Ge3 and Ge corresponding to an ohmic contact does not provide a conduction channel to charge carriers. The experimental Seff of the Mn5Ge3/Ge bilayer can be predicted independently of layer thicknesses using the model of parallel conduction channels along the temperature gradient . ΔV was measured across the Mn5Ge3 step while the Ge side of the sample was located either in the colder part or the hotter part  or the hot region (orange open squares) of the temperature gradient. The Seebeck coefficient is different from the single films of Ge and Mn5Ge3, as well as from that measured on the Mn5Ge3/Ge bilayer  or from Mn5Ge3-to-Ge (cold Ge), showing clearly an impact of the interface on Seff. Seff exhibits higher values if the carriers move from hot Mn5Ge3 to cold Ge, and depending on the carrier direction through the interface, Seff follows opposite variations versus temperature (increasing or decreasing versus temperature), exhibiting almost a symmetrical behavior around the value of ~ 30 µV K−1. Figure 5Ge3-step/Ge-terrace sample. The measurements are reproducible. However, few differences in the Seebeck values are observed between same-condition measurements  SGe in the case of the Mn5Ge3/Ge serial-phase geometry. However, the bipolarity of the Ge film should be considered. Indeed, the Ge film being extrinsic, it contains two different charge carriers of significantly different concentrations: holes (h) and electrons (e), being respectively the majority and the minority charge carriers in the present case. Thus, a similar model as for the parallel-phase geometry should be used to consider the two conduction channels of holes and electrons in Ge, giving:Figure Sh and Se the respective Seebeck coefficients of holes and electrons, and σh and σe their respective conductivities in the Ge film (σGe = σh + σe). Typically, the coefficients α and (1−α) respectively correspond to the volume fractions of the two phases Mn5Ge3 and Ge in the considered bi-phase material. Though these coefficients should describe the volume difference of Mn5Ge3 and Ge crossed by the charge carriers between the two electrodes placed on the two sides of the Mn5Ge3 step considering the measurement geometry shown in Fig. Seff depends on (1) the position of the electrodes during measurement  the variations of S5/3 versus temperature, and (3) the Ge bipolarity versus temperature, without considering any interface effect.With ts α and −α respec69. Furthermore, the n-type Mn5Ge3/p-type Ge interface corresponds to an ohmic contact  or a source (carrier injection) of charge carriers in the Mn5Ge3 and Ge films depending on the carrier flux direction towards the Mn5Ge3/Ge interface. The Mn5Ge3/Ge interface is expected to locally modify the concentration of the charge carriers, and thus, to have a combined effect on the five parameters S5/3, Sh, Se, σh, and σe in Eq.  and the electron concentration (n) in the semiconductor, one can roughly assume Sh ∝ p−2/3, Se ∝ n−2/344, σh ∝ p, and σe ∝ n55, leading to the approximationThe interface is an extended defect that can hold interfacial carrier traps, as well as carrier recombination centersact Fig. b, and she in Eq. , leadingp increases from 1013 to 1018 cm−3 while n = 5 × 1015 cm−3 (orange solid line) or if n increases from 1013 to 1018 cm−3 while p = 5 × 1015 cm−3 (blue solid line). The plot uses two arbitrary constant A and B according to S = A + B (p1/3 − n1/3)/(p + n), which were adjusted in order to obtain Seebeck coefficient variations of the same order as those in Fig. p and n variations  can be interpreted as an increase of the fraction n/p versus T, while Seff variations when the charge carriers cross the interface from Ge-to-Mn5Ge3 (hot Ge) can be interpreted as a decrease of the fraction n/p versus T.Figure 5Ge3-step/Ge-terrace sample while the temperature gradient was oriented along the Mn5Ge3 step. A first measurement was performed on the Mn5Ge3 step, and a second was performed on the Ge terrace. The first measurement should be similar to Seff measured on Mn5Ge3/Ge bilayers  and Ge (p-type) films, a strong electron and hole concentration gradient should be present at the interface vicinity  in Fig. 5Ge3 being about one order of magnitude higher than the hole concentration in Ge according to Hall effect measurements  after deposition allowing for the growth of the Mn5Ge3 layer via reactive diffusion and the full crystallization of the Ge layer70. The thickness of the Mn layer was maintained the same for all the samples, while the thickness of the Ge film was varied in order to obtain a 160 nm-thick Mn5Ge3 layer on a Ge layer of different thicknesses. Hall measurements and sample resistivity were measured in the Van der Pauw geometry using a lab-made setup operating between 20 and 350 K. The applied magnetic field for Hall measurements was 0.5 T. The Seebeck coefficients of the films were measured using a home-made setup71 between T = 225 and 325 K. The distance d between the two electrodes allowing to simultaneously determine the potential difference ΔV as well as the temperature gradient ΔT = Th − Tc was 1 cm (Fig. The Ge and Mncm2 Fig. a glass s


Microbial CH4 production is the result of a combination of many different microorganisms and microbial processes, which together achieve the mineralization of organic matter to CO2 and CH4. Temperature dependence applies to each individual step and each individual microbe. This review will discuss the different aspects of temperature dependence including temperature affecting the kinetics and thermodynamics of the various microbial processes, affecting the pathways of organic matter degradation and CH4 production, and affecting the composition of the microbial communities involved. For example, it was found that increasing temperature results in a change of the methanogenic pathway with increasing contribution from mainly acetate to mainly H2/CO2 as immediate CH4 precursor, and with replacement of aceticlastic methanogenic archaea by thermophilic syntrophic acetate-oxidizing bacteria plus thermophilic hydrogenotrophic methanogenic archaea. This shift is consistent with reaction energetics, but it is not obligatory, since high temperature environments exist in which acetate is consumed by thermophilic aceticlastic archaea. Many studies have shown that CH4 production rates increase with temperature displaying a temperature optimum and a characteristic apparent activation energy (aE). Interestingly, CH4 release from defined microbial cultures, from environmental samples and from wetland field sites all show similar aE values around 100 kJ mol−1 indicating that CH4 production rates are limited by the methanogenic archaea rather than by hydrolysis of organic matter. Hence, the final rather than the initial step controls the methanogenic degradation of organic matter, which apparently is rarely in steady state.There is virtually no environmental process that is not dependent on temperature. This includes the microbial processes that result in the production of CH There is virtually no microbial activity that would not be regulated by temperature . BiologiMicrobial life is complex. Microorganisms contain many different enzymes, which may react differently upon temperature changes. Microbial populations consist of many different individual microorganisms, each possibly slightly different in its response to temperature. In nature, microbial communities consist of many different microbial populations with different physiologies and life styles and thus, with potentially different features of temperature dependence. Therefore, it is by principle very complex how environmental microbial communities will react to temperature changes and it is hard to make any predictions.Nevertheless, it is worthwhile to review the literature on temperature dependence of microbial activity to see whether there are any guiding principles. In the following I will primarily (but not exclusively) focus on methanogenic microbial communities living in anoxic environments such as flooded rice fields and aquatic sediments, whose temperature characteristics have frequently been studied over the last 40–50 years. A previous mini-review on this subject  will be 2 + CO2, formate, trimethylamine, dimethylsulfide, and methanol to CH4. The most common methanogenic substrates in nature are acetate and H2 + CO2 . The mK value is the substrate concentration at maxV/2. In this way, metabolic rates are modulated by the substrate availability. Both maxV and mK of enzymatic reactions are expected to increase with increasing temperature (maxμ S/(SK + S), where SK is the substrate concentration at maxμ/2. Note, however, that mK and SK are different values with a different physiological meaning and relevance.Methane is produced by methanogenic archaea, which convert simple substrates like acetate, HH2 + CO2 . The ratperature . Microbi4 production increase with increasing temperature. The increase in rate (v) can be described by the Arrhenius equation, i.e., v = A exp {−aE/(RT)} with A = Arrhenius constant and aE = apparent activation energy. In a microbial cell, e.g., a methanogenic archaeon, CH4 is produced as the end product of a catabolic reaction chain, in which the methanogenic substrate is converted to CH4  face a cell-specific affinity that increases with temperature, whereas substrates without active transport  face a temperature independent cell-specific affinity , such as predicted by 4 production rates are not limited by the substrate transport of substrates, but by one of the subsequent enzymatic reaction steps. This limiting enzyme reaction may be the same as that defining the kinetic isotope effect of CH4 production (Methanosarcina versus Methanothrix (Methanosaeta) species, respectively . These data are from a meta-analysis of literature data, based on 33 different strains of growing or non-growing methanogenic archaea  during steady state than during the phase of excess substrate  is close to that of defined methanogenic populations (106 kJ mol−1) suggesting that CH4 production by methanogenic archaea is the rate-limiting step also when organic matter is the primary substrate (2 + CO2) when temperature increases. An alternative explanation is that the hydrolysis of organic matter, as the initial rate-limiting step under steady state conditions, has a similar aE value as the CH4 production by methanogenic archaea  shows a narrow range of kJ mol−1 . The meta values . Neverthpiration , photosypiration  or to thanalysis .aE values of methanogenic archaeal populations with those of entire ecosystems, one should realize that at the level of ecosystems the temperature changes are recorded as daily rhythms and seasonal variations. In addition to the temperature effects, ecosystems are exposed to many other control variables, such as variations in substrate supply [e.g., root exudations in vegetated soil and sediment .Nevertheless, microbial and enzymatic activity is also found at the low eratures . This caeratures . Howevereratures . This iseratures , indicatBacteria . For exaembranes , thus afembranes . In facto, ΔHo, and ΔSo for reactions important in CH4 production pathways can be calculated from tabulated values , which is −39.94 kJ mol−1 protons produced. Temperature dependence of ΔGo is basically a function of the magnitude of the process entropies (ΔSo) under standard conditions. The actual ΔG at the given concentrations of reactants and products can be calculated from the Nernst equation, which is also temperature-dependent, i.e., ΔG = ΔGo + RT ln(P/S), where P are the products and S the reactants. P and S represent the product of the activities or partial pressures of the individual products and reactants, respectively. At dilute solutions activities may be replaced by concentrations. The lowest possible concentrations of the reactants are those when the actual ΔG = 0, i.e., the threshold. For example, hydrogenotrophic methanogenesis proceeds until the threshold concentration  of H2 is reached. Below the threshold H2 can no longer be consumed and CH4 production stops }) with the Michaelis–Menten equation  change of a reaction is given by the changes of the enthalpy and the entropy, i.e., ΔG = ΔH – TΔS. For standard conditions  values of ΔGes e.g., ; Table 1on stops . On the  acetate . The degntly low . The thry ΔG = 0  and thosit lower . The dev–1/3 ATP . The H2 equation , or by iequation .2 partial pressures are a function of the standard free enthalpies of anaerobic H2-consuming processes, and generally decrease with the useful energy (more negative ΔGo). Hence, the H2 threshold is for example lower in hydrogenotrophic ferric iron reducers < sulfate reducers < methanogens < homoacetogens. This has been found in microbial populations  and 33% (1 CH4) to total CH4 production (3 CH4), respectively. The ‘syntrophic’ pathway is quite common for many methanogenic environments including anaerobic digestors (The pathway starts with primary fermentation (1) of sugar to lactic acid, followed by secondary syntrophic fermentation (2) of lactic acid to acetate, COigestors , rice fiigestors , peat boigestors  and lakeigestors , especiain-situ temperature (Many studies have shown that the anaerobic degradation pathway can change with the perature . At low perature :6H12O6 → 2CH3CHOHCOOH1: C3CHOHCOOH + 2H2O → 2CH3COOH + 2CO2 + 4H22: 2CH2 + 2CO2 → CH3COOH + 2H2O5: 4H3COOH → 3CH4 + 3CO24: 3CH6H12O2 → 3CH4 + 3CO2Sum: COr it may happen when sugar fermentation to lactic acid (1) is replaced by homoacetogenic sugar fermentation (6):6H12O6 → 3CH3COOH6: C3COOH → 3CH4 + 3CO24: CH6H12O2 → 3CH4 + 3CO2Sum: C4 is not produced from acetate, although it is the dominant product of organic matter degradation. Instead, it is only degraded by oxic respiration and ferric iron reduction (2 + CO2 is not excluded.It has been shown for many anoxic environments that homoacetogenesis is a favored process at low (<15°C) temperatures , so thateduction . Hence, 4, respectively, as observed in anaerobic digestors  coupled to hydrogenotrophic methanogenesis (3) (Also at high temperature (>45°C), aceticlastic and hydrogenotrophic methanogenesis can operate producing 67 and 33% of the CHigestors  and in rigestors . This isigestors . Frequenesis (3) :6H12O6 → 2CH3CHOHCOOH1: C3CHOHCOOH + 2H2O → 2CH3COOH + 2CO2 + 4H22: 2CH3COOH + 4H2O → 8H2 + 4CO27: 2CH2 + 3CO2 → 3CH4 + 6H2O3: 12H6H12O2 → 3CH4 + 3CO2Sum: CThe existence of syntrophic acetate oxidizing bacteria (7) and their general preference for high temperatures has been shown in anaerobic sludge digestors  and rice2/CO2), low temperatures (psychrophilic) by the acetogenic pathway (<67% acetate), and high temperatures (thermophilic) by the hydrogenotrophic pathway (100% H2/CO2), such preferences are by no means obligatory. The mechanistic reasons for these preferences are not quite clear. A look at the thermodynamics of the degradation pathways may be helpful. The values of ΔGo are constrained by the stoichiometries of the pathways. The values change with temperature according to the magnitude of ΔSo  is always −405, −426, and −445 kJ at 5, 25, and 45°C, thus moderately decreasing (becoming more exergonic) with temperature. The ΔGo’T values of the individual reactions at 5, 25, and 45°C are shown in 2 become increasingly less endergonic or more exergonic, if temperature increases. o, 2 + CO2 is thermodynamically less favorable than hydrogenotrophic methanogenesis, and results in relatively higher H2 thresholds  are dominated by the ‘syntrophic’ pathway . Sharing reduces the length of the pathway (number of enzymatic steps) for the individual species, which may be favorable for energetic reasons  and aceticlastic methanogenesis both become thermodynamically more favorable when temperature increases . Thus, wosarcina . Anotherosarcina  or more  reasons . Since t2, acetate and CO2 as final products, (2) hydrogenotrophic methanogens and (3) aceticlastic methanogens , in which energy is usually shared by at least three different microbes, (1) the secondary fermenters producing Hhanogens . Three pMethanosarcina and Methanothrix species and an increase of thermophilic Methanocella species (formerly RC-I), which are hydrogenotrophic  are especially abundant at low temperatures. These taxa have also been found in a boreal mire, but with little changes in community composition over a seasonal temperature gradient between 0 and 14°C . Interestingly, methanogenic archaea were found to exhibit a rather high range of aE values of >100 kJ mol−1. In the environment, hydrolysis of complex organic matter is the first step of the methanogenic degradation processes, and this process can exhibit markedly lower aE values of about 60 kJ mol−1. However, analysis of various methanogenic environments with complex microbial communities display aE values, which are not similar to those of hydrolysis but to those of the methanogenic archaea. This observation indicates that hydrolysis of organic matter is not the rate limiting step of CH4 production in most environments, meaning that the microbial community is frequently supplied with pulses of easily degradable substrates.Microorganisms proliferate within and tolerate a more or less wide range of temperatures. This is observed in environments that experience only small temperature fluctuations  and also in others that experience dramatic temperature changes on a daily or seasonal range . The microorganisms generally display characteristic temperature optima and a characteristic increase of reaction kinetics with increasing temperature, which can be modelled by the Arrhenius equation using a characteristic apparent activation energy (The author confirms being the sole contributor of this work and has approved it for publication.


In this article by Xi Xiang et al.,Figure 4In Figure 4A, the HE staining of Day 1 in control group and Day 7 in BMSCs + MB(SDF‐1α)(IA) + US group were exactly the same. The corrected Figure 4A is below.Figure 5In Figure 5, the reddish O‐fast green staining of Day 2 in control group and Day 1 in control group were exactly the same. Day 2 control group was mistakenly placed in Day 7 control group. The corrected Figure 5 is below.The authors confirmed that all results and conclusions of this article remain unchanged.We apologise for these errors.