As rapid progress has been achieved in emerging thin film solar cell technology, organic–inorganic hybrid perovskite solar cells (PVSCs) have aroused many concerns with several desired properties for photovoltaic applications, including large absorption coefficients, excellent carrier mobility, long charge carrier diffusion lengths, low‐cost, and unbelievable progress. Power conversion efficiencies increased from 3.8% in 2009 up to the current world record of 22.1%. However, poor long‐term stability of PVSCs limits the future commercial application. Here, the degradation mechanisms for unstable perovskite materials and their corresponding solar cells are discussed. The strategies for enhancing the stability of perovskite materials and PVSCs are also summarized. This review is expected to provide helpful insights for further enhancing the stability of perovskite materials and PVSCs in this exciting field. Park and co‐workers further significantly enhanced the PCE to 6.5% in 2011.3−xClx perovskite deposited on the mesoporous Al2O3 film that was deposited on a compact TiO2 layer to demonstrate a PCE of 10.9% upon contacting 2,2′,7,7′‐tetrakis‐9,9′‐spirobifluorene (spiro‐MeOTAD) as hole transport layer (HTL), which implied that perovskite materials can transport the electrons by itself without electron transporting materials.2 achieved 15.0% in 2015,2 device was achieved to 19.6%.FigureAlthough PVSCs have been intensively studied in solar cells recently, these materials were first discovered by Weber 35 years ago. The 3D structure of perovskite materials was applied by Mitzi to transistors and light‐emitting diodes.The development of PVSCs has focused on two important issues with respect to practical applications. One is photoelectric conversion efficiency and the other is stability. The PCE of PVSCs has been increased from 3.1 to 22.1% in just a few years through materials design and device engineering based on the desired optoelectronic properties, while the PVSCs are now more competitive than those of traditional crystalline solar cells as a promising photovoltaic technology. Hence, the real challenge lies on the long‐term stability of PVSCs. PVSCs are prone to decay when exposed to humidity, heat, light, oxygen, and so on. At the same time, it is also important to explore device components, device structures, and interface properties. To be marketable within the industry, PVSCs must be able to operate continuously for close to 25 years in outdoor conditions. Long‐term stability tests for PVSCs without encapsulation should be over 500 h under 1 sun illumination23). The general formula for the organic–inorganic halide perovskites is 3D ABX3, where A is CH3NH+ (MA+) or NH = CHNH+ (FA), Cs+, B is a metal ion , X can be I−, Br−, or Cl−.64− octahedral, and A cation is located in the cavity between four BX64− octahedra and X is surrounded by 12 nearest neighbors n−1(CH3NH2)nPbI3][H3O], subsequently reducing the stability of perovskite.3 undergoes serious degradation by the introduction of water molecules into the crystal structure, where they form weak hydrogen bonds to the highly hygroscopic methylammonium cations leading to bond dissociations between the crystal constituents.2 and perovskite.Therefore, moisture degradation of perovskite is related with the MAPbI3A stable perovskite layer with high quality is crucial for the long‐term stability of PVSCs. This is because perovskite layer is very sensitive to external environment, consequently leading to the degradation of the PVSCs. In general, the stability of the perovskite layer depends on its chemical components, crystal structure, and the film quality of perovskites.3.13 can be determined by a tolerance factor t and an octahedral factor µ, t = (RA + RX)/{√2 (RB + RX)}, µ = RB/RX, where RA, RB, and RX are the ionic radii of the corresponding ions. Generally, to form a stable perovskite structure at room temperature, the value of t is between 0.81 and 1.11. When t is 0.9–1.0, it is expected to be ideal cubic perovskite, while distorted structures orthorhombic, rhombohedral, tetragonal, or hexagonal are likely to be formed when t is larger or smaller. At present, the value of t usually is 0.834 (R(CH3NH3+) = 0.18 nm, R(Pb2+) = 0.119 nm, and R(I−) = 0.220 nm) to the CH3NH3PbI3 perovskite. The t can be adjusted by the replacement or partial introduction of ions with different sizes to further obtain more stable cubic crystal structure.The stability of PVSCs crystal structure ABX3.1.12)2+) were replaced to MA+, t value increased to 0.99, which is higher than t = 0.91 for CH3NH3PbI3  perovskite. The larger FA iron can also expand the lattice6 octahedra,3 to ≈1.49 eV for FAPbI3.6 octahedra was stronger than that of MA cation due to higher probability to form hydrogen bonding.3 under light was reported to the formation of hydroiodic acid (HI) by the released protons from MA cation with I−. However, protons released for FA cation is less than those released from MA cation due to resonance characteristics of FA cation via C—N bonds.3 than for MAPbI3. Compared to MAPbI3, FAPbI3 is not sensitive to high temperature. While MAPbI3 discolors in 30 min, FAPbI3 does not discolor even at 150 °C under ambient conditions.3 is found to show poorer humidity stability compared to MAPbI3.3 (α‐phase) was reported to be turned to yellow nonperovskite phase (β‐phase) in the presence of liquid interface at room temperature. Furthermore, FA+ cation tended to decompose into ammonia, which further degraded the moisture‐stability of FAPbI3.The organic cation (A) is a key part of the perovskite and can determine the perovskite crystal structure and its stability. When the slightly larger formamidinium ion FA , the degree of degradation was severer for FAPbI3 film (85.9%) than for FA0.9Cs0.1PbI3 film (65.0%) after 19 h.0.9Cs0.1PbI3 film was much better than the pristine FAPbI3 film under 85% relative humidity at 25 °C in dark condition, due to more stabilized FA ions. The t value of Cs+ incorporated FA1−xCsxPbI3 perovskite was systematically investigated by Zhu and co‐workers.0.85Cs0.15PbI3 with 15% Cs+ has a t value of 0.9 with a stable cubic phase. When FAPbI3 and FA0.85Cs0.15PbI3 thin films were exposed to a humid environment for 18 d, the absorbance value of FAPbI3 thin films decreased very fast while that of FA0.85Cs0.15PbI3 thin films remained constant. The PCE of the FA0.85Cs0.15PbI3 PVSCs showed almost no change while that of FAPbI3 decreased to about half the initial value when exposed to humid air for 400 h. Moreover, Park et al. also reported Cs/FA mixtures perovskite films enhanced the thermal and humidity stability of PVSCs, reaching a PCE of 16.5%. The improved structural stability was explained by that Cs was effective in assisting the crystallization of the black phase in FA perovskite due to entropic stabilization.x(MA0.17FA0.83)(1−x)Pb(I0.83Br0.17)3 with Cs/MA/FA triple cation achieved PCE of 21.1% and output of 18% after 250 h under operational conditions.As presented in 3 is also critical for the structural stability of the perovskite materials.2+) with a smaller volume to replace Pb2+ caused a decline in the stability of the lattice.2+ is rapidly oxidized to more stable Sn4+ which is harmful to the structural stability of the perovskite.0.75Cs0.25Sn0.5Pb0.5I3 with a wider‐band gap FA0.83Cs0.17Pb(I0.5Br0.5)3 perovskite materials has also been investigated as potential systems with the PCE of monolithic two‐terminal tandem PVSCs achieving 17.0%. Remarkably, such infrared‐absorbing perovskite cells exhibited excellent thermal and atmospheric stability. Obviously, we need further fundamental studies on the performance of Pb‐free perovskite and more advanced solar cell preparation techniques to develop high‐performance Pb‐free PVSCs in the future.Apart from A cation, the metal B atom in ABX3 is 2000 h, which does not meet the requirement of the practical applications and commercialization. Incorporation of a small atomic radius of Br or Cl with similar properties into MAPbI3 perovskite has been found to be an effective approach to improve the device stability. Noh et al. tuned the stoichiometry of MAPb(I1−xBrx)3 perovskites by substituting I− with Br−.1−xBrx)3  hybrid solar cells suffered from serious degradation after exposure to 55% humidity, whereas the PCE of perovskite with higher Br concentration maintained stable with no obvious degradation observed after 20 d. A low sensitivity to humidity of the cells based on higher Br content was associated with a reduced lattice parameter and a transition from tetragonal to cubic phase. When Br− ions were introduced into the perovskite structure, the lattice parameters were changed, for instance, 5.921 for the MAPbBr3, 6.144 for the MABr/MAI 2:1, and 6.223 for the MABr/MAI 1:2. The change in the lattice parameter is mainly on account of the difference between the ionic radii of Br− (1.96 Å) and I− (2.2 Å).− is smaller, which is also favorable for the formation of the cubic structure. Misra et al. reported the substitution of larger I atoms with smaller Cl atoms in MAPbI3−xClx and stability of PVSCs improved. The authors showed that the improved stability can be related to its compact and stable structure, which led to a reduction of the lattice constant and a transition to a more stable cubic phase.3−yCly PVSCs with improved stability. The devices retained 95% of its initial PCE after storage in glovebox without any device encapsulation for 30 d.1−xBrx)3−yCly PVSCs with different Br content, the PVSCs were stored in a glovebox under nitrogen atmosphere in the dark. The MAPbI3−yCly PVSCs without Br took a PCE drop of 20% in 30 d, however with 50% Br content there was a notable increase of about 37% in PCE, due to the rearrangement of 3D perovskite.The highest reported stability of PVSCs with light absorbing MAPbIt value of perovskite structures.−) is a stable pseudohalogen anion with similar properties to I− ion. Liang and co‐workers fabricated MAPbI3−x(SCN)x with a small amount of Pb(SCN)2 adding to a PbI2 solution. The partial substitution of I− by SCN− in the perovskite structure could remarkably improve the stability compared to the conventional iodide‐based perovskite.2I perovskite by substituting I− in MAPbI3 with two SCN−.3 and MAPb(SCN)2I thin films exposed to 95% relative humidity air for several hours was investigated in the reflection spectrum. A substantially increased reflection value for MAPbI3 in 0.5 h indicated MAPbI3 perovskite quickly degraded under humidity. In contrast, the MAPb(SCN)2I reflection value increased slightly under the same humidity for 4 h. Moreover, the color of the MAPbI3 film was almost completely changed to yellow in air for 30 d, indicating the formation of PbI2, whereas no significant change was observed for MAPb(SCN)2I. When exposed to 95% relative humidity air, the PCE of MAPbI3‐based device decreased from 8.8 to 6.9% in 7 d and completely decomposed in 14 d, while that of the MAPb(SCN)2I‐based device only decreased from 8.3 to 7.4% in 14 d. Yan and co‐workers also prepared a stable MAPbI3−x(SCN)x in humidity air.− groups into the perovskite lattice was thermodynamically stable. Moreover, strong ionic interactions existed between SCN− and the adjacent Pb2+ ions as well as hydrogen bonds forming between SCN− and MA+, which helped to improve chemical stability.With similar chemical characteristics and ion radius to halogen anions, pseudohalogen anions can also affect the 3.1.26H5(CH2)2NH3+ (PEA+),+,3(CH2)3NH3+ (BA+)3 lattice. As presented in Table 2(MA)2[Pb3I10] (PEA = C6H5(CH2)2NH3+, MA = CH3NH3+) for the first time by integrating PEA+ into the 3D MAPbI3 perovskite.3 and (PEA)2(MA)2[Pb3I10] are compared in Figure2(MA)2[Pb3I10] perovskite films can be deposited through spin‐coating without thermal annealing. To compare the moisture stability of (PEA)2(MA)2[Pb3I10] and MAPbI3, both materials were exposed to 52% relative humidity. The (PEA)2(MA)2[Pb3I10] perovskite films kept the same X‐ray diffraction (XRD) patterns after the exposure for 46 d 2(MA)2[Pb3I10], resulting in lower device performance than 3D perovskite.In recent years, the application of 2D layered perovskite materials has been considered as a promising possibility material to improve the stability of PVSCs. The 2D perovskite materials were prepared by the insertion of larger cations, for example, C2, MAI, and PEAI to yield compounds with different layer (n) values in the series PEA2(MA)n−1PbnI3n+1.n = ∞ corresponds to the cubic 3D perovskite (CH3NH3PbI3) while other n values describe 2D (n = 1) or quasi‐2D (n > 1) perovskite structures  perovskites. As a result, the device stability was remarkably improved. The PCE of the 3D perovskite dropped to 18% of its original PCE after eight weeks when stored in N2, nevertheless devices based on the quasi‐2D perovskites kept 74% (n = 60) and 86% (n = 40) of its initial PCE after the same time under a low‐humidity atmosphere. It was concluded that quasi‐2D perovskites possessed much higher formation energies as a result of the appreciable van der Waals forces that confer to improve stability compared to 3D perovskites. The additional organic cations with larger radius can actually protect the perovskite structure as a capping layer in quasi‐2D perovskite.PEA‐based layered perovskite films were investigated by continuously tuning the dimensionality of perovskite compounds with mixing different stoichiometric quantities of PbI3(CH2)3NH3)2(MA)n−1PbnI3n+1 (n = 1–4) thin films was reported by employing a large alkylammonium ion CH3(CH2)3NH3+(BA+) n−1PbnI3n+1 perovskites showed better moisture resistance compared to their 3D MAPbI3 analogues. The authors ascribed the moisture‐resistance of the (BA)2(MA)n−1PbnI3n+1 to the hydrophobicity of the long BA+ chain and the highly oriented and dense nature of the perovskite films, which prevented the direct contact of water molecules with the perovskite.The fabrication of the 2D lead iodide perovskite  in γ‐butyrolactone (GBL), DMF, or dimethyl sulfoxide (DMSO) to a clear solution.Figure3 perovskite film by one‐step deposition with the antisolvent engineering. In FDC procedure, a perovskite precursor solution is spin‐coated on the TiO2 layer, after a specific delay time, following by that a second solvent (chlorobenzene) is quickly added to the substrate to achieve the high‐quality perovskite thin films. The role of the second solvent is to rapidly reduce the solubility of CH3NH3PbI3 in the solvent and thereby promoting fast nucleation and growth of the crystals in the film  can be used as a template for the preferential crystal growth of MAPbI3 in the mesoporous oxide host by forming hydrogen bonds between its COOH and NH3+ groups and I− ions from the PbI6 octahedra. The COOH group of 5‐AVA coordinated with mesoporous TiO2 and ZrO2 and the NH3+ ions on the terminal of 5‐AVA+ served as nucleation sites for the formation of the perovskite crystal 5‐AVA+.+ modified perovskite worked stably for more than 1000 h in humid air and under full sunlight. Nazeeruddin and co‐workers reported 1,1,1‐trifluoro‐ethyl ammonium iodide (FEAI) additive in MAPbI3 perovskite to show a better morphology and greatly enhance the stability of the perovskite.3NH3PbI3 perovskite via strong hydrogen bonding of the —PO(OH)2 and —NH3+.−2 ultraviolet‐filtered simulated sunlight at 45 °C for one week with performance of CH3NH3PbI3‐ABPA retaining 90%.Chu and co‐workers investigated alkali metal halides  into the perovskite precursor solution as additives to obtain significant improvements in the perovskite morphology.Xu and co‐workers fabricated highly smooth and complete coverage perovskite films by involved hydrochloride to perovskite films process with large crystalline grains. The authors proposed that the formation of Pb—I—Pb is much more difficult when chloride is coordinated to lead, leading to a slow nucleation and growth. When hydrochloride gas was dissolved in the perovskite precursor solution, the coordination tended to form chloride coordinated ions due to the high affinity between lead and chloride. Therefore, addition of hydrochloride enabled the improvement of perovskite film morphology with a large grain size. Moreover, the devices were highly air‐stable and the 95% of the initial PCE (16.9% certified) remained after 30 d storage in air at room temperature without encapsulation.44.13NH3PbI3 was employed as the dye sensitizers in dye‐sensitized solar cells (DSSCs)  Figure6. Ther4.1.12, a mesoporous metal oxide TiO2/Al2O3 as the transparent n‐type component filled with perovskites, followed deposited by perovskite layer, a HTL of spiro‐MeOTAD, and a metal anode (Au or Ag)  cathode [fluorine doped tin oxide (FTO)], a compact electron transport layer (ETL) TiO) Figure . Kim and4.1.22/perovskite/spiro‐OMeTAD/Au. In 2013, Ball et al. reported the first planer PVSCs with the PCE of 12.3%.3NH3PbI3−xClx solar cells achieved an improved PCE of 15.4% by vapor‐phase deposition.VOC and JSC relative to a mesoscopic device processed with the same materials, the planar device usually holds more severe J–V hysteresis.Compared with the mesoporous structure of PVSCs, planar heterojunction solar cells have increased the flexibility optimization, providing the possibility for the development of the tandem solar cells, large‐scale industrial production, and the further study of the working mechanism for PVSCs. Planar heterojunction perovskite solar cells are usually composed of FTO conductive glass, compact barrier layer, perovskite layer, and hole conductor layer, which can be further divided into planar p‐i‐n and planar n‐i‐p structure. Figure 61BM) or ‐phenyl‐C71‐butyric acid methyl ester (PC71BM). The perovskite devices with the p‐i‐n structure provided the advantages of high efficiencies, lower temperature processing, flexibility, and negligible J–V hysteresis effects. The first planar p‐i‐n structure of PVSCs was delivered by Guo and co‐workers in 2013.x and TiO2 layers have recently been used for the hole and electron selective contacts, respectively, which made the perovskite device distinct from its organic counterpart. Yang and co‐workers reported a solution‐processed lead halide perovskite solar cell with a maximum value of 16.1% based on p‐type NiOx and n‐type ZnO nanoparticles as HTL and ETL, respectively.For another planar structure with the p‐i‐n configuration, the HTL is on top of the transparent conducting substrate, as shown in Figure 2 scaffold structure can promote collection of photogenerated charge carriers and enhance perovskite potential for future photovoltaics.In terms of the two kinds of device structures, ultimately all devices may adopt the simple structure of planar PVSCs configuration, representing a potentially important step in the path toward commercialization. However, at present, mesoscopic structures reached a certified stable PCE of 22.1%, which still holding a substantial lead over planar structures. The mesoporous TiO4.22 has been used as the most frequently efficient inorganic ETLs in most PVSCs, owing to its favorable energy level, easy fabrication, and long electron lifetimes.2 increase the electron recombination rates, reducing the efficiency and stability under UV illumination of the device. Further results indicated that charge accumulation induced by ion migration in the perovskite layer could be the main cause of hysteresis.FigureTable2.3NH3PbI3 when contact with ZnO.2 have higher mobility than TiO2 and better stability with a band gap of above 3.6 eV, making it conceptually a more likely candidate in high‐performance PVSCs.2 nanoparticles have successfully been utilized as the ETL for PVSCs with a PCE of 20.5% with almost free of hysteresis.3 (LBSO) with high electrical mobility was also employed as perovskite ETL.The ETL can be used to protect the perovskite layer out of degradation suffering from external environments. TiO2‐based devices suffered from the degradation upon UV illumination, doping is an effective way to modify the electrical properties of TiO2 and improve the stability of PVSCs. Han and co‐workers successfully fabricated large‐size (>1 cm2) PVSCs with n‐doped (n+) TiOx matrix with Nb5+ ions to improve rapid carrier extraction by avoiding pinholes and cracks over large areas. The result revealed that PVSCs maintained 97% of the initial PCE after 1000 h of light soaking with the Ti(Nb)Ox layer protecting the perovskite from intrusion of humidity.2 could remove oxygen defects from the TiO2 lattice, leading to the improvement of the performance and enhancement in stability.2 is also helpful to improve the device stability upon UV light. Ito et al. inserted a Sb2S3 layer between the TiO2 and CH3NH3PbI3 interface to block the UV‐induced photocatalysis and reduce the degradation of perovskite. While the control device almost did not work after 12 h exposure to light, with TiO2/Sb2S3 as ETL, PCE still maintained 65% of the initial value after the 12 h. This result again confirmed that UV‐induced reaction at the TiO2/CH3NH3PbI3 interface was partly responsible for the degradation of the perovskite.Since the TiO61BM, ICBA, and PC71BM.60 to MAPbIxCl3−x and FA0.83Cs0.17Pb(I0.6Br0.4)3 PVSCs. The 4‐‐N,N‐diphenylaniline (N‐DPBI) dopant provided a higher conductivity and tuned the surface wettability of the C60 films. The devices achieved a PCE of 18.3% and long‐term stability owing to the reduced number of surface defects and filling of trap states with higher electron density.Fullerene and its derivatives are the most widely used n‐type materials for organic ETLs, such as PC4.3tert‐butylpyridine (TBP) is an essential step to enhance the hole mobility and suppress charge recombination.2.3h Truxene core with high mobility and suitable surface energy. As an effective HTL without the need of extra doping process, the excellent PV performance was achieved with a PCE of 18.6% and better stability.The HTL is also vital to prepare highly efficient and stable PVSCs. As a whole, the HTL can be divided into four families: small molecule, polymeric, inorganic, and carbon.3),3NH3PbI3 film with large crystalline grains on PTAA HTL which dramatically reduced charge trap density and enhanced the PCE to ≈18.1% for planar PVSCs with its higher hole mobility.Compared to 3D structural spiro‐MeOTAD, conducting polymer materials own higher hole mobility and film‐forming property, constituting a major class of organic HTL in PVSCs. PEDOT:PSS is the generally adopted HTL.x,x,x as HTL in the inverted planar p‐i‐n PVSCs,x film suggested superior performance than that of device with PEDOT:PSS layer. Kamat and co‐workers reported another copper‐based inorganic p‐type hole‐conductor CuI as the p‐type HTL in mesoscopic PVSCs,3NH3PbI3 films were fabricated with the CuSCN layer as the HTL,3−xClx (CuSCN) films hold lower surface roughness and smaller interface contact resistance between the perovskite layer and the selective contacts, exhibiting PCE of over 18%. Furthermore, the PVSCs based on CuSCN presented higher reproducibility and stability.Inorganic materials were explored to be good candidates as HTL owing to their intrinsically properties, such as NiOCarbon‐based HTL offers one of the more stable solutions even though the device efficiency seems to be not yet comparable to the organic HTLs. Seok and co‐workers developed reduced graphene oxide (RGO) at the first time to enhance conductivity and the stability of the PVSCs.4.42.Metal electrodes are an important part in PVSCs for charge collection. The common metal electrodes are silver, aluminum, and gold in the PVSCs.4.53NH3PbI3−xClx perovskite devices with a mesoporous Al2O3 scaffold were encapsulated with an epoxy resin and a glass coverslip in a nitrogen‐filled glove box.For commercial solar cells, such as silicon solar cells, CIGS solar cells, etc., encapsulation methods are applied to improve the stability of the devices.5Although the PCE exceeding 22% has been achieved in PVSCs with the unprecedented rapid development, commercialization of PVSCs will require the synergistic development of high‐efficiency, large‐area with long‐term stability as a competitive technology. However, the actual long‐term stability of PVSCs still lags behind their outstanding efficiency. Thus, the various causes of devices degradation need to be further explored in‐depth to achieve high performance and stable PVSCs. The stability of perovskite layer is usually optimized by compositional engineering of halides and cations, the application of 2D materials and additives. On the other hand, incorporating proper charge transport layers and electrodes or interfacial resistant agents into PVSCs is also considered as effective ways to maintain device stability as well as the efficiency of PVSCs.xBr3−x perovskite material is more appropriate for the high stability of 3D perovskite layer. In addition, 2D systems provide new interesting challenges and the possibilities to new physical properties, also regarding charge transport, making them currently gain more attention in the field. Apart from the perovskite materials, stable device structures, interfacial engineering to form a stable and efficient interface interaction, inert electrode and proper encapsulation are also crucial. To date, the PVSCs with the best stability were achieved with mesoporous structure of FTO/c‐TiO2/meso‐TiO2/perovskite/PTAA/Au. In our view, the further improved mesoporous device structure has a great practical application potential due to the simplicity and ease of fabrication with decent performance. However, deposition of the mesoporous layer by avoiding high‐temperature annealing process would make more convenient for fabricating large‐area and the long‐term stability of PVSCs. The commonly used TiO2 is sensitive to UV with presence of surface defects (oxygen vacancies) in TiO2, which leads to the degradation of PVSCs. Therefore, exploring stable novel ETLs to replace TiO2 is highly desirable. For instance, the solution‐processed SnO2 is the good candidate for the high‐performance electronic devices instead of TiO2. Furthermore, with regard to the HTLs, developing doping‐free, cheap organic material instead of expected Li salt‐doped spiro‐MeOTAD or developing the HTL‐free structure devices is very attractive. For the planar p‐i‐n PVSCs, some stable and efficient inorganic, such as NiOx HTL is potential to replace PEDOT:PSS. Hence, fabricating PVSCs with reliable high efficiency and long‐term stability is a systematic and very multidisciplinary project, which requires contributions from several directions.Therefore, the ultimate challenge for PVSCs is to achieve stable devices and perovskite materials that can survive in an outdoor environment. For the perovskite layer, ions selection, doping, and crystal structure are promising to improve the perovskite layer stability. For example, FA/MAPbIMuch progress has been already made in the stability of the halide perovskite. Nevertheless, the underlying mechanism of the degradation remains far from being comprehensively understood. To realize the industrialization development of PVSCs, an in‐depth comparative understanding of the degradation mechanisms is essential for establishing efficient solutions to improve the PCE and the stability of PVSCs.The authors declare no conflict of interest.


The authors would like to report a typographical error that states the concentration of staurosporine used for the induction of apoptosis is 10 μM. The actual concentration used is 1 μM. The authors confirm that the changes do not affect the results of the experiments.The typographical error can be found in the following two instances and are corrected here:In the Figure 8 legend, Panel F (page 17 in the PDF):“Primary cultures were pre-treated with KA (1 mM) for 2 h prior to adding apoptotic NE-4C cells (treated with 5 μM CM-DiI for 25 min and 1 μM staurosporine for 4 h).”Under Materials and Methods, in the second paragraph under the ‘Primary Microglia Cultures’ heading (Page 32 in the PDF):“NE-4C were previously labeled with the membrane marker CM-DiI  and treated with staurosporine  to induce apoptosis.”


Spherical materials with yolk‐shell structure have great potential for a wide range of applications. The main advantage of the yolk‐shell geometry is the possibility of introducing different chemical or physical properties within a single particle. Here, a one‐step hydrothermal synthesis route for fabricating amphoteric yolk‐shell structured aluminum phenylphosphonate microspheres using urea as the precipitant is proposed. The resulting microspheres display 3D sphere‐in‐sphere architecture with anionic core and cationic shell. The controllable synthesis of aluminum phosphates with various morphologies is also demonstrated. The anionic core and cationic shell of the aluminum phenylphosphonate microspheres provide docking sites for selective adsorption of both cationic methylene blue and anionic binuclear cobalt phthalocyanine ammonium sulphonate. These new adsorbents can be used for simultaneous capture of both cations and anions from a solution, which make them very attractive for various applications such as environmental remediation of contaminated water. Making use of the gradual hydrolysis of urea in the synthesis system, the morphology of the resultant aluminum phenylphosphonates could be modulated by varying the reaction time.22.16H7O3P) and aluminum nitrate to form yolk‐shell structured aluminum phenylphosphonate microspheres. Urea has been selected as a precipitating reagent because its hydrolysis is much slower, which in contrast to the precipitants such as NaOH and NH3·H2O, is well suited for precise control of the pH value needed for monitoring the morphology evolution. The morphology of the as‐synthesized YSNs material (ys‐AlPhPO) after 3 h of hydrothermal treatment (HT) was established by scanning electron microscopy . The sample exhibits type IV adsorption–desorption isotherm with H3‐type hysteresis loop, which indicate the presence of mesopores, echoing with HRSEM results that the non‐uniform mesopores are formed by the disorderly assembled rod‐like nanoparticles. The Brunauer–Emmett–Teller (BET) specific surface area and total pore volume of the sample are calculated to be 163 m2 g−1 and 0.28 cm3 g−1, respectively. The powder X‐ray diffraction (XRD) pattern of ys‐AlPhPO  exhibits a single strong diffraction peak at 2θ of 5.8°. The energy‐dispersive X‐ray spectroscopy (EDX) result of ys‐AlPhPO suggests that the Al/P ratio is 0.81. The inductively coupled plasma (ICP) and elemental analyses of ys‐AlPhPO indicate that the content of Al, P, C, and H is 9.20%, 12.78%, 33.03%, and 3.58%, respectively. The calculated Al/P ratio of 0.83 is in a good agreement with the result of the EDX analysis. As expected, the 13C magic angle spinning (MAS) nuclear magnetic resonance (NMR) spectrum shows peaks in the chemical shift range from 125 to 135 ppm representing resonance signals of the benzene ring . The solid‐state NMR together with the FT‐IR data  demonstrate the incorporation of phenylphosphonate into the framework through the coordination between aluminum and phenylphosphonate, and confirm that the phenylphosphonate can retain the structural integrity during the hydrothermal synthesis process.The porous structure of ys‐AlPhPO microspheres is further evidenced by N2.227Al MAS NMR, 31P MAS NMR, and powder XRD as illustrated in FigureTo investigate the formation mechanism of the hierarchical yolk‐shell structured aluminum phenylphosphonate microspheres through one‐step precipitation and hydrothermal synthesis process, a series of samples was synthesized by varying hydrothermal treatment time. The crystallization processes and morphological evolution as a function of HT time were monitored by SEM, 27Al MAS NMR spectrum of cs‐AlPhPO synthesized for 2 h of HT exhibits two resonance peaks centered at 42.5 and −22.4 ppm 3⋅H2O, which is similar to that in the previous report.[[qv: 8d]] Further prolongation of the HT time to 2.5 causes an increase in the pH value from 1.6 to 2.1 in the synthesis system, which results in the conversion of the flower‐like microspheres to the smooth‐surfaced microspheres through an Oswald ripening. When the HT time reaches 3 h and the pH value increases to 4.6, the outside surface of the smooth‐surfaced microspheres is transformed to the mesoporous shell through an Oswald ripening and chemical structure variation, meanwhile the unconverted smooth‐surfaced microspheres act as a yolk to form yolk‐shell structures through electrostatic interactions. This can be evidenced by the fact that the smooth‐surfaced microspheres serve as the yolk with a negative potential of −4.1 mV and the spiny‐type microspheres resemble shells with a positive potential of +24.6 mV at pH 4.6. Increasing HT time to 4 h causes an increase in the pH value to 7.3 and the hydrolysis of aluminum forms Al(OH)4− and Al(OH)3 in the reaction solution. At pH exceeding 5, the yolk is converted to the shell by consumption of Al(OH)4− or Al(OH)3 leading to the formation of spiny‐type microspheres with hollow structures. Afterward, the shell of hollow‐structured microspheres grows inward and the cavity becomes thick with prolonged HT time. The morphological transformation of AlPhPO microspheres from flower‐like to smooth‐surfaced and yolk‐shell, and then to hollow structures is accompanied by the evolution of the chemical structure from Al2(O3PC6H5)3⋅H2O to Al4(O3PC6H5)5(OH)2⋅nH2O, and Al5(O3PC6H5)6(OH)3⋅nH2O, and finally to Al7(O3PC6H5)3(OH)15⋅nH2O as speculated on the basis of elemental analysis, solid NMR, and powder XRD data. Briefly, the morphological transformation of microspheres with prolonged HT time can be ascribed to: (1) structural change, more specifically, the change of aluminum coordination environment derived from tetra‐coordination aluminum further coordinated to hydroxyl with pH value increasing in the synthesis system, (2) Oswald ripening process, and (3) CO2 and NH3 derived from urea after prolonging the HT time serve as a bubble‐type template for formation of hollow space and porous structures. The morphological transformation from core‐shell to yolk‐shell and next to hollow spheres has also been observed in previous studies for different compositions.The plot of the zeta potential versus pH  indicates that the electrostatic interactions between the cationic shell and the anionic yolk control the formation of these smooth‐faced microspheres. The formation mechanism is illustrated in 2.3Very interestingly, these anionic cores and cationic shells of yolk‐shell structured aluminum phenylphosphonate microspheres (ys‐AlPhPO) can effectively uptake both cationic methylene blue (MB) and anionic binuclear cobalt phthalocyanine ammonium sulphonate (PDS) as shown in Figure Figure−1. The equilibrium adsorption data were analyzed using two adsorption models, Langmuir and Freundlich models, in order to establish the most appropriate adsorption isotherm. The linear plots made for ss‐DNA on ys‐AlPhPO are shown in insets of Figure The promising application of these microspheres has been demonstrated for bio‐adsorption and enrichment; namely, ys‐AlPhPO microspheres exhibited an excellent performance for adsorption of salmon sperm DNA Figure6A and −9m from the MS spectrum; whereas the enriched phosphopeptides from the digests of the casein protein can be identified successfully after exposure to the microspheres. The above results indicate the strong potential of microspheres as promising alternatives for enrichment of low abundance peptides.The microspheres were also employed to enrich low concentration of biomolecules for matrix‐assisted laser desorption‐ionization time‐of‐flight mass spectrometry (MALDI TOF MS) detection. Alpha‐conotoxin IMI, a functional peptide (sequence GCCSDPRCAWRC) was used as model, which can selectively block alpha‐delta site of the muscle acetylcholine receptor. The signal strength of alpha‐conotoxin IMI can be raised to over 45 000  with the same morphology as their precursors, aluminum phenylphosphonate microspheres, can be obtained by calcination at 550 °C under air atmosphere for 2 h. No visible change in the particle size and structure was observed from SEM images Figure7. At a327Al MAS NMR. The amphoteric yolk‐shell structured microspheres are formed by electrostatic interactions between the anionic yolk and cationic shell, and can be used for selective adsorption of dyes. The yolk‐shell structure and unique amphoteric properties make aluminum phenylphosphonate microspheres suitable for various applications in many fields such as environment, catalysis, biology, and nanomedicine.In summary, a one‐step precipitation and hydrothermal method have been developed for the synthesis of amphoteric yolk‐shell structured aluminum phenylphosphonate microspheres with mesoporous shells. Aluminum phenylphosphonate microspheres with different surface morphologies, from flower‐like to smooth‐surfaced microspheres and yolk‐shell microspheres, and then to hollow structured microspheres, can be facilely obtained by varying the hydrothermal treatment time. The observed morphological transformation together with the change in the composition and structure of aluminum phenylphosphonate microspheres with reaction time can be attributed to the variation of Al coordination, evidenced by 4Chemicals and Reagents: All materials were of analytical grade and used as received without any further purification. PPA , aluminum nitrate nonahydrate, (Al(NO3)3⋅9H2O, ≥98%), α‐cyano‐4‐hydroxycinnamic acid , trifluoroacetic acid , acetonitrile , and proteins/peptides were purchased from Sigma‐Aldrich. Salmon sperm DNA for adsorption and release was obtained from Wako Pure Chemical Industries Ltd. and used as received . Other reagents were purchased from ShangHai Chemical Reagent Inc. of Chinese Medicine Group.Synthetic Procedure: x‐AlPhPO microspheres were synthesized via a hydrothermal method using C6H7O3P and Al(NO3)3⋅9H2O in the presence of urea as the precipitation reagent. In a typical synthesis, C6H7O3P (0.3 mmol) and Al(NO3)3⋅9H2O (0.3 mmol) were dissolved into 15 mL of HNO3 solution (0.035 mol L−1) and then urea (15 mmol) was added into the above solution. The mixture was stirred for 30 min until a homogeneous solution was formed; next, the resultant solution was transferred into a Teflon‐lined autoclave and heated at 100 °C for n h  under static conditions. After cooling to room temperature naturally, the solid products were collected using a membrane filter (0.22 μm) and washed with distilled water repeatedly, and dried at room temperature. The as‐synthesized materials were denoted as x‐AlPhPO, where x = cs, ys, and hs, refer to the core‐shell, yolk‐shell, and hollow structures, respectively.Characterization: PXRD patterns were recorded on an Empyrean powder diffraction system using Cu Kα radiation of 0.15406 nm wavelength. The nitrogen adsorption experiments were performed at −196 °C on a Micromeritics ASAP 3000 system. Prior to the measurement, the samples were outgassed at 120 °C for at least 6 h. The BET specific surface areas were calculated using adsorption data in the relative pressure range of P/P0 = 0.05–0.25. Pore size distributions were calculated from the adsorption branch using the Barrett–Joyner–Halenda (BJH) method. The total pore volumes were estimated from the amounts adsorbed at a relative pressure (P/P0) of 0.99. HRSEM, field‐emission scanning electron microscopy (FESEM), and SEM images were taken on a HITACHI S‐5500, HITACHI SU 8010, and JSM 6360 LV microscope operating at an accelerating voltage of 1–30 kV, respectively. TEM imaging was performed using a JEOL JEM‐2010 at an acceleration voltage of 120 kV. FT‐IR spectra were collected with a TENSOR 27 IR spectrometer in the range of 4000–400 cm−1 using KBr pellet. 13C (100.5 MHz) cross‐polarization magic angle spinning (CP‐MAS), 31P (161.8 MHz), and 27Al (79.4 MHz) MAS solid‐state NMR experiments were recorded on a BRUKER DRX 400 spectrometer equipped with a magic angle spin probe in a 4 mm ZrO2 rotor. 13C signals were referenced to tetramethylsilane, 31P NMR signal was referenced to H3PO4 (85 wt%). The experimental parameters were 6 kHz spin rate, 2 s pulse delay, and 6 min contact time for 13C CP‐MAS NMR experiments; 8 kHz spin rate, 3 s pulse delay, and 10 min contact time for 31P MAS NMR experiments.Adsorption of Cationic MB and Anionic Binuclear Cobalt PDS: cs‐AlPhPO, ys‐AlPhPO, and hs‐AlPhPO were dried at 80 °C for 24 h in a vacuum oven prior to adsorption of dyes. In each adsorption experiment, 100 mg of x‐AlPhPO were added to 3 mL of MB (0.01 mg mL−1) and PDS (0.05 mg mL−1) solutions, and the resulting mixture was continuously shaken in a shaking bath at room temperature for 12 h and then was standing for 12 h.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


We show that the rigid hollow cavities of the CNT enforce size-controlled nanoparticles which are electrochemically active inside the CNT. The ferrimagnetic Mn3O4 filling is switched by electrochemical conversion reaction to antiferromagnetic MnO. The conversion reaction is further exploited for electrochemical energy storage. Our studies confirm that the theoretical reversible capacity of the Mn3O4 filling is fully accessible. Upon reversible cycling, the Mn3O4@CNT nanocomposite reaches a maximum discharge capacity of 461 mA h g−1 at 100 mA g−1 with a capacity retention of 90% after 50 cycles. We attribute the good cycling stability to the hybrid nature of the nanocomposite: (1) Carbon encasements ensure electrical contact to the active material by forming a stable conductive network which is unaffected by potential cracks of the encapsulate. (2) The CNT shells resist strong volume changes of the encapsulate in response to electrochemical cycling, which in conventional  Mn3O4 hinders the application in energy storage devices. Our results demonstrate that Mn3O4 nanostructures can be successfully grown inside CNT and the resulting nanocomposite can be reversibly converted and exploited for lithium-ion batteries.The ferrimagnetic and high-capacity electrode material Mn Manganese features a plethora of stable oxidation states and thus forms various oxides, such as Mn(II)O, Mn2(III)O3, Mn3O4, and Mn(IV)O2. A promising theoretical specific capacity of 937 mA h g−1 in case of Mn3O4, which is nearly three times higher than that of graphite (372 mA h g−1)2, has raised considerable interest in Mn3O4 as anode material6. However, the strong fading of the electrochemical capacity due to fractionation, resulting from pronounced volume changes associated with the conversion reaction, as well as low electric conductivity seriously hinder its applicability in secondary batteries7. Nanosizing promises enhanced capability to accommodate strain, which is induced by electrochemical cycling, and may reduce kinetic limitations of the macroscopic counterparts of electrode materials9. Consequently, various manganese oxide/carbon based hybrid nanomaterials have been reported to at least partly solve these issues. These hybrid materials show enhanced electrochemical properties, such as high specific capacities and good cycling stability14. Carbon nanotube (CNT) based composites include MnO/CNT mixtures15 as well as the exohedral functionalization of CNT by Mn3O42. So far, however, neither the synthesis of Mn3O4-filled CNT nor electrochemical studies on this material have been reported, as previous studies solely addressed manganese oxide particles located on the outer surface of the nanotubes17. Owing to their excellent conductivity, chemical stability, and very high mechanical strength, CNT as such have indeed proven to be a promising choice as carbon source in hybrid nanomaterials18. However, in conventional approaches using exohedrally functionalized CNT the synthesis of uniformly sized and shape-controlled nanoparticles of different manganese oxides is an unsolved challenge. The approach presented here applies the wet chemical synthesis of crystalline Mn3O4 nanoparticles inside the inner hollow cavities of multi-walled carbon nanotubes  via a solution-based approach.Downsizing well-established materials to the nanoscale is a key route towards novel functionalities, in particular if different functionalities are merged in hybrid nanomaterials. One example are manganese oxides which show a relatively low electromotive force, high natural abundance, and environmental benignity and hence are considered as one of the best candidates for an anode material in lithium-ion batteries (LIB)3O4@CNT hybrid nanocomposites are obtained after filling CNT with a manganese salt solution and a subsequent reducing step (cf. Refs20). With a reduction profile of 4 h at 500 °C under Ar/H2 flow, homogeneously MnO-filled CNT (MnO@CNT) are synthesized. The XRD-pattern of MnO@CNT in comparison with reference Bragg peak positions (ICSD #16203921) confirms the presence of phase-pure MnO with space group Fm3O4@CNT, as confirmed by the XRD pattern in Fig. 3O4 phase with space group I41/amd (ICSD #6817422). Again, all additional reflections are ascribed to the CNT. The XRD data hence suggests the complete oxidation of MnO to Mn3O4 upon thermal treatment. Note, that all patterns show relatively broad Bragg reflections, typical for the small size of the MnO and Mn3O4 nanocrystals.Mnns ICSD #6203921 c3O4-filled CNT. In both materials, manganese oxide nanoparticles are located mainly inside the CNT. From the TGA data , which is smaller than the size-limiting inner diameter of the utilized CNT (~35 nm). The broad size distribution of the oxidic particles inside the CNT reflects that, in addition to the particles featuring pearl necklace-like structures, MnO- and Mn3O4@CNT exhibit smaller nanoparticles , both in the range of 0.01–3.0 V vs. Li/Li+. Figure st, 2nd, and 10th cycle of a cyclic voltammogram (CV), recorded at a scan rate of 0.1 mV s−1. During the initial cycle, starting with the cathodic scan, five distinct reduction peaks (I-V) and three oxidation peaks (i-iii) are observed. The redox pair I/i around 0.1 V and the irreversible reduction peak III at 0.7 V can be attributed to processes related to multi-walled CNT, as supported by the CV of unfilled CNT which shows the same features  on the surface of the CNT23. The pronounced redox pair I/i indicates intercalation and deintercalation of Li+-ions between the graphitic layers of the CNT25. In this regard, the slight splitting of the oxidation peak i resembles the staging phenomenon reported for graphite electrodes26. The observed increase of this peak intensity (i) during cycling is also found in the CV of the unfilled CNT  is reflected by the reduction peak V at 1.3 V vs. Li/Li+. The adjacent shoulder around 1.45 V may be associated with the reduction of small amounts of amorphous manganese oxides with a higher Mn oxidation state than that in Mn3O44, e.g., Mn5O829. Proceeding to lower potentials, peak IV at 0.9 V is associated with the reduction of LiMn3O4 to MnO (B); and the pronounced reduction peak II corresponds to the conversion of MnO to metallic Mn (C). The associated oxidation from Mn back to MnO is reflected by peak ii at 1.3 V during the first anodic scan. The additional oxidative feature iii around 2.1 V might display the back-formation of Mn3O431. Interestingly, this feature occurs in only few studies31, which might point to enhanced reaction kinetics promoted by the conductive CNT network3. In the second cycle, the reportedly irreversible reduction peaks IV and V have nearly vanished, while peak II splits into a double peak and shifts to approx. 0.4 V. The shift indicates a structural transformation due to conversion reaction (C) where disordered Mn and Li2O are formed accompanied by a volume expansion33. There are no significant changes between cycle 2 and 10, which indicates the good cycling stability of the Mn3O4@CNT nanocomposite.During the first cathodic scan Fig. , the int3O4@CNT filling as well as the presence of intermediate products appearing upon electrochemical cycling is also confirmed by magnetization studies, as such experiments have been reported to be an appropriate tool to confirm electrochemical conversion in transition metal oxides36. Indeed, the magnetic properties of the associated manganese oxides differ strongly so that large magnetization changes are expected. The conversion reactions (A) and (B) reflect the switching of ferrimagnetic Mn3O4 nanoparticles to antiferromagnetic MnO. The respective magnetic ordering phenomena appear at TC = 42 K in Mn3O4 and TN = 120 K in MnO39. Switching of the magnetic properties upon electrochemical treatment is directly confirmed by the magnetization data in Fig. 3O4@CNT indeed shows a long-range magnetic order at TC = 42 ± 1 K, confirming the ferrimagnetic nature of the encapsulate. Upon galvanostatic lithiation at 5 mA g−1 down to 0.5 V, i.e. after passing the reduction peaks V, IV, and III  and  reflect 3O4@CNT in comparison to the pristine material (a) are displayed in Fig. 3O4 during initial lithiation, which probably yields an agglomeration of several nanoparticles inside single CNT. The volume expansion is also apparent from a lower density of the encapsulate, which can be deduced from a lower contrast to the CNT environment in the TEM image. No clear differences are observed between the lithiated  and delithiated (c) Mnted Fig.  and the ted Fig.  material3O4@CNT is further studied by means of charge and discharge profiles of a rate capability test with specific charge/discharge currents between 50 and 1000 mA g−1  vanish, most probably due to the overpotential caused by the ohmic resistance of the cell setup. In addition, the features corresponding to the conversion reaction (C) (II/ii) become more sloping but stay present.The electrochemical performance of Mng−1 Fig. . In the  CV Fig.  can be i3O4@CNT and pristine CNT at 100 mA g−1 (GCPL) are compared in Fig. −1, respectively. The CNT reach a maximum discharge capacity of 334 mA h g−1 in the 2nd cycle which decreases moderately to 299 mA h g−1 in cycle 50. The Mn3O4@CNT composite exhibits increasing capacities for approximately 15 cycles before reaching a maximum discharge capacity of 463 mA h g−1 in cycle 18, of which 93% is maintained after 50 cycles (429 mA h g−1). Thus, the incorporation of Mn3O4 into CNT leads to more than 40% enhanced specific capacities on average as compared to unfilled CNT. The initial capacity increase was also observed in other studies on Mn3O4/CNT composites11 and could originate from Mn3O4 nanoparticles which are not properly attached to the conductive CNT network before electrochemical cycling.The progression of charge/discharge capacities of Mn3O4 (29.5 wt%) to the theoretical value associated with the reversible conversion reaction step (C), i.e. 703 mA h g−1, the data have been corrected for the contribution of the CNT and normalized accordingly , and then begin to decline with a capacity retention of around 90% after 50 cycles. The contributed capacities even exceed the theoretical expectations of the conversion reaction (C) from the 6th cycle on. This observation can be explained by an additional contribution due to oxidative feature iii  and 64% (500 mA g−1). Subsequently, they increase simultaneously to the specific capacities and demonstrate satisfying values between 98–99% from the 8th cycle on. Those values indicate a decent cycling stability of the Mn3O4@CNT composite. Further measurements regarding the cycling stability at 100 mA g−1 confirm that the depicted decline stays very moderate with a capacity loss of approx. 15% per 50 cycles for a total of more than 100 cycles. The capacity losses can be attributed to different reasons: on the one hand, the small amount of Mn3O4 nanoparticles located on the outside of the CNT may upon volume expansion either detach or inhibit Li+ transfer from/into the CNT; on the other hand the CNT themselves do not offer perfect cycling stability  filling the open CNT with a metal salt solution, (2) reducing the metal salt to MnO, and (3) oxidation of MnO to Mn3O4. For the first step, a solution of Mn(NO3)2·3H2O  with a concentration of 1 mol/L and CNT  were dispersed for 1.5 hours in an ultrasonic bath. This homogeneous dispersion was filtered and the tubes were dried. In the second step, the dry, filled CNT were reduced in an Ar/H2 flow (50 sccm min−1/50 sccm min−1) for 4 h at 500 °C. Subsequently, the material was tempered at 350 °C for 6 h in Ar flow (50 sccm min−1) to form Mn3O4@CNT.The Mnα1 radiation (λ = 1.5406 Å), scanning electron microscopy  and transmission electron microscopy . SEM images were obtained either in the back scattered electrons (BSE) mode or in the secondary electrons (SE) mode. Measurements of the particle diameter were accomplished with the program ImageJ40. A SDT Q600 (TA Instruments) was used for thermogravimetric analysis (TGA). During TGA measurements the filled CNT were burned at a heating rate of 5 K min−1 up to 850 °C under the flow of synthetic air. Magnetic measurements of Mn3O4@CNT were performed by means of a MPMS-XL5 (Quantum Design) SQUID magnetometer with powder samples. The temperature was varied between 2 and 300 K according to zero-field-cooling (ZFC)/field-cooled-cooling (FCC) procedures at 100 Oe. Hysteresis loops were obtained at 5 and 300 K in magnetic fields of up to ± 5 T.The composite was characterized by X-ray diffraction ) using Cu K3O4@CNT composite as well as pristine CNT were characterized electrochemically by means of cyclic voltammetry (CV) and galvanostatic cycling (GCPL). The measurements were performed in Swagelok-type two electrode cells on a VMP3 potentiostat (BioLogic) at a constant temperature of 25 °C. The working electrodes were prepared by stirring the active material in a solution of polyvinylidene fluoride  in N-methyl-2-pyrrolidone  overnight, and then evaporating most of the NMP in order to obtain a slurry which was spread on circular copper current collectors. The weight ratio of active material to PVDF amounted to 86:14 and the mass loading of the electrodes was 3–4.5 mg cm−2. Afterwards, the electrodes were dried at ~100 °C in vacuum (<5 mbar) overnight, mechanically pressed at 10 MPa, and then dried again. The cells were assembled in an argon atmosphere glove box (O2/H2O < 1 ppm), using a lithium metal foil counter electrode pressed on a nickel current collector, two layers of glass microfibre separator (Whatman GF/D), and 200 μl of a 1 mol l−1 solution of LiPF6 in 1:1 ethylene carbonate and dimethyl carbonate (Merck Electrolyte LP30).The MnThe datasets generated during and/or analysed during the current study are available from the corresponding author on reasonable request.Supplementary Material S1-S3


TPH2) catalyses the initial and rate‐limiting step in the biosynthesis of serotonin, which is associated with a variety of disorders such as depression, obsessive compulsive disorder, and schizophrenia. Full‐length TPH2 is poorly characterized due to low purification quantities caused by its inherent instability. Three truncated variants of human TPH2  were expressed, purified, and examined for changes in transition temperature, inactivation rate, and oligomeric state. chTPH2 displayed 14‐ and 11‐fold higher half‐lives compared to rchTPH2 and NΔ47‐rchTPH2, respectively. Differential scanning calorimetry experiments demonstrated that this is caused by premature unfolding of the less stable regulatory domain. By differential scanning fluorimetry, the unfolding transitions of rchTPH2 and NΔ47‐rchTPH2 are found to shift from polyphasic to apparent two‐state by the addition of l‐Trp or l‐Phe. Analytical gel filtration revealed that rchTPH2 and NΔ47‐rchTPH2 reside in a monomer–dimer equilibrium which is significantly shifted toward dimer in the presence of l‐Phe. The dimerizing effect induced by l‐Phe is accompanied by a stabilizing effect, which resulted in a threefold increase in half‐lives of rchTPH2 and NΔ47‐rchTPH2. Addition of l‐Phe to the purification buffer significantly increases the purification yields, which will facilitate characterization of hTPH2.Tryptophan hydroxylase 2 ( TPH uses the cofactor Fepterin BH to gener‐HT. TPH hTPH, the catalytic cores of both isoforms have been structurally and enzymatically characterized Along with phenylalanine hydroxylase (PAH) and tyrosine hydroxylase (TH), TPH forms an enzyme subfamily of iron(II)‐containing mono‐oxygenases collectively referred to as the aromatic amino acid hydroxylases (AAAHs). Various truncated variants of all three AAAHs have been structurally characterized All enzymes in the AAAH family form tetramers through dimers of dimers Escherichia coli expression systems rnPAH)  which comprises the catalytic and the regulatory domains. This structure lacks interpretable electron density in the first 18 residues of the N terminus, suggesting that this region is flexible rnPAH and hTPH2 , it is found that residues 20 to 47 of hTPH2 align with the mobile 18 residues of rnPAH, which suggests that these residues of hTPH2 are mobile and might cause instability and insolubility. Therefore, an NΔ47‐rchTPH2 variant was expressed, purified, and characterized.The available crystal structures of TPH1 only comprise the catalytic domain, while the crystal structure of TPH2 includes both the catalytic and the tetramerization domains hTPH2 , rchTPH2 , and NΔ47‐rchTPH2 . The tetramerization domain (residue 460–490) was removed to investigate the role of the regulatory domain when TPH2 was not in a tetramer. As TPH2 is poorly characterized due to low stability, additional efforts were made to identify ligands that could increase the stability, and hence purification yield, of TPH2.The current study sheds light on the influence of the regulatory domain on macromolecular structure and stability of TPH2 by characterizing three truncated variants (sequences shown in Supporting information); cl‐phenylalanine, l‐tryptophan, and l‐tyrosine), other ligands known to bind certain ACT domains . For chTPH2, an almost ideal two‐state unfolding behavior was observed, Fig. Tm value of chTPH2, Fig. hTPH2 from two‐state to continuous; hence, no Tm values could be obtained at high ligand concentration.Very little is known about the regulation and structure of TPH as purification of this enzyme results in limited quantities. The presence of the regulatory domain is known to cause low stability and solubility  range 0. μm–10 mmhTPH2, gave rise to inconsistent and polyphasic unfolding curves from which no Tm values could be obtained . The ligand screen showed that l‐Trp and l‐Phe induced unfolding transitions with apparent two‐state unfolding in a concentration‐dependent manner . The same stabilizing trend for l‐Phe and l‐Trp was observed for NΔ47‐rchTPH2, Fig. Containing the regulatory domain, rcl‐Trp and l‐Phe gave rise to increased transition temperatures, a combination of the compounds was analyzed. Figure l‐Trp or l‐Phe in the presence of the other compound results in increased Tm values, suggesting that l‐Trp and l‐Phe are able to increase the transition temperature in an additive fashion. The additive effect is supported by the increase in Tm values with increasing l‐Phe concentration observed at saturated concentration of l‐Trp . In consensus with the findings of D'Sa et al. hTPH2 was found to elute at a volume corresponding to the molecular weight of a monomer (36.2 kDa) m, chTPH2 was found only to reside as a monomer, Fig. m l‐Phe did not induce any change in the elution pattern of the monodisperse solution complementing the observations from the DSF experiments.It has previously been shown in DSF assays that under stabilizing conditions multi‐component complexes can change from polyphasic to almost two‐state unfolding and that this is indicative of monodispersity and stability et al. hTPH2, as SEC of NΔ47‐rchTPH2 yielded two overlapping peaks with elution volumes corresponding to molecular weights of a monomer (47.4 kDa) and a dimer (94.9 kDa)  Fig. . The graequation . A Kd vhTPH2 and rchTPH2/NΔ47‐rchTPH2 shows that dimerization is caused by the regulatory domain. This finding is supported by a model structure proposed by Jiang et al. et al. et al. The difference in oligomeric states of cl‐Phe on the monomer–dimer equilibrium was analyzed by performing SEC on NΔ47‐rchTPH2 samples with and without the addition of l‐Phe. In Fig. hTPH2 without the addition of l‐Phe, and the orange curves represent the elution of protein when the sample and running buffer contained 3 mm l‐Phe. It is evident that l‐Phe shifts the equilibrium toward dimer. At a loading concentration of 2 μm NΔ47‐rchTPH2, the addition of l‐Phe was found to shift the equilibrium from predominately monomer to almost exclusively dimer. The same shift in monomer–dimer equilibrium was observed for rchTPH2 , the monomer–dimer equilibrium does not seem to be influenced by 1.0 mm l‐Trp as monomer is present  of the hTPH2 variants were calculated using t½ = ln(2)/k. At 30 °C and a concentration of 5 μm, rchTPH2, and NΔ47‐rchTPH2 displayed t½ values of only 15 ± 2 and 18 ± 3 min, respectively. Hence, truncation of the N terminus slightly increased the t½ value . With a t½ value of 203 ± 40 min, chTPH2 displayed a 14‐fold higher half‐life compared to rchTPH2, which is in agreement with the results of Carkaci‐Salli et al. m l‐Phe, t½ values of 41 ± 3, 49 ± 5, and 193 ± 29 min were obtained for rchTPH2, NΔ47‐rchTPH2, and chTPH2, respectively. Variants containing the regulatory domain displayed nearly 3‐fold increase in half‐lives whereas chTPH2 was not significantly influenced by l‐Phe.To investigate if the dimerizing effect had an impact on TPH2 stability, rates of inactivation at 30 °C in the presence or absence of 3 m °C Fig. . From thhTPH2 and NΔ47‐rchTPH2 implies that l‐Phe is stabilizing through binding to the regulatory domain. Further, as l‐Phe induces a stabilizing effect in both rchTPH2 and NΔ47‐rchTPH2, binding of l‐Phe does not seem to involve the 47 N‐terminal residues. l‐Phe induced stability and shift in equilibrium toward dimer imply that TPH2 variants containing the regulatory domain are more stable as dimers.That the stabilizing effect is exclusively observed for rcl‐Trp has previously been observed by McKinney et al. l‐Trp might stabilize through binding in the active site, which might tighten up flexible regions, and hereby protect the active site. The increase in Tm (DSF) upon the addition of l‐Phe or l‐Trp might, analogously to the hypothesis of McKinney et al. l‐Phe, suggesting that the increase in Tm and half‐life occur through different mechanisms. Alternatively, l‐Phe binds in an allosteric site which stabilizes the regulatory domain through improved interactions with an adjacent monomer. Such an allosteric site in the regulatory domain has been identified in PAH, which has been found to increase PAH activity and stability and induce large conformational changes l‐Trp and l‐Phe observed in the DSF experiments. Such an allosteric site might have relevance in vivo functioning as an allosteric modulating site that stabilizes a dimer in the native tetramer (dimer of dimers). Stabilization of the tetramer is important as disruption of the native tetramer results in decreased enzymatic activity A stabilizing effect of Tm, of the TPH2 variants were measured utilizing differential scanning calorimetry (DSC), to investigate if the low half‐life of rchTPH2 is caused by premature unfolding of the regulatory domain. The DSC experiments were performed in the presence of l‐Phe to observe domain unfolding in a monodisperse solution. All variants displayed irreversible unfolding upon reheating and hence, only Tm values were extracted from the thermograms. Illustrative examples of DSC thermograms of the three variants are presented in Fig. The transition temperatures, hTPH2 and NΔ47‐rchTPH2 ). This suggests that the lowest temperature transitions are partly due to the unfolding of the regulatory domain, as suggested for hPAH by Thórólfsson et al. Tm(high)). Transition temperatures obtained from the thermograms are presented in Table Tm values of the two main transitions by 1.6 and 2.9 °C, respectively. Truncation of the N terminus only induced a slight increase in the Tm(high) value .The unfolding transitions of the three variants seem equivalent, however, by comparison of the thermograms of cPH2 Fig. A,B, it ihTPH2 unfold at 40.1 ± 1.2 °C and 46.8 ± 0.3 °C, respectively. These results are, despite variations in the buffer systems, in the same range as previously reported transition temperatures for AAAHs (47.5–55.5 °C) hTPH2  compared to chTPH2 (12.0 ± 2.4 mg·L−1). The low quantities for rchTPH2 were overcome by truncation of the N terminus which resulted in an 11‐fold increase in yield (10.4 ± 3.1 mg·L−1), which was similar to that of chTPH2.Characterization of l‐Phe was reflected in the purification yields, as they were increased seven‐ and twofold for rchTPH2 (6.4 ± 1.2 mg·L−1) and NΔ47‐rchTPH2 (25.1 ± 2.9 mg·L−1), respectively .The stabilizing effect of ely Fig. . In linel‐Trp and l‐Phe change the unfolding mechanism of hTPH2, but only when the regulatory domain is present. Deletion of the C‐terminal tetramerization domain results in a monomer–dimer equilibrium which is shifted to predominately dimer with the addition of l‐Phe. In the presence of l‐Phe, the dimer displayed significantly increased half‐life which in turn resulted in significantly increased purification yields of hTPH2 variants containing the regulatory domain. These findings will facilitate future characterization of hTPH2.The current results demonstrate that All used chemicals were of analytical grade, and all solutions were prepared using water from an 18.2 MΩ·cm Milli‐Q synthesis A10 Q‐Gard system which was filtered through a 0.22‐μm filter. Protein purification was performed on an ÄKTA purifier 100 from GE Healthcare. Utilized GE Healthcare column variants: HiLoad Superdex 200 26/60 pg, Superdex 200 10/300 GL, and a XK 16/20 column packed with 25 mL Dextrin Sepharose High Performance media. During purifications, all TPH‐containing solutions were kept in ice water, except during the chromatographic steps, which were performed at room temperature. Protein solutions were concentrated using an Amicon ultrafiltration cell with an Ultracell PL‐3 membrane. Protein concentrations were determined by measuring the absorbance at 280 nm on an ND‐1000 NanoDrop Spectrophotometer from Saveen Werner .E. coli was obtained from GenScript . All proteins were expressed as maltose binding protein fusion proteins from the pET26 expression vector in E. coli BL21(DE3)  cells. The sequences of the proteins expressed are given in the Supporting information. The recombinant fusion proteins contain a cleavage recognition site for human rhinovirus 3C protease (3CP) Full‐length human TPH2 cDNA optimized for expression in Escherichia coli cells from 650 mL cultures were thawed from −80 °C and resuspended in buffer containing 20 mm HEPES/NH4OH, 300 mm (NH4)2SO4, pH 7.0  or 20 mm HEPES/NH4OH, 300 mm (NH4)2SO4, pH 7.0, and 3 mm l‐Phe (Phe buffer), to a volume of 40 mL. (NH4)2SO4 was included in the purification buffer, as it has been found that the purification yield of recombinant catalytic core of rabbit TPH1 was enhanced by the addition of ammonium sulfate g for 20 min. The supernatant was decanted to another tube and centrifuged a second time at 4 °C and 18 000 g for 20 min, while the pellet was discarded. The supernatant was collected and filtered through a 0.45‐μm filter. A volume of approximately 35 mL filtered supernatant was loaded with a flow rate of 5 mL·min−1 onto a Dextrin Sepharose column, which was equilibrated with five column volumes of Std or Phe buffer. Following sample loading, an MBP‐3CP solution was prepared by diluting MBP‐3CP from stock with Std or Phe buffer to a concentration of 2.2 μm. About 30‐mL MBP‐3CP solution was loaded onto the column with a flow rate of 5 mL·min−1, and the column was incubated at room temperature for 1 h. Once the flow (5 mL·min−1) was resumed, 10 mL of the protein‐containing eluate was collected. The collected solution, containing target protein, was filtered through a 0.45‐μm filter prior to loading on a HiLoad Superdex 200 prep grade column, which had been equilibrated with two column volumes of Std or Phe buffer. Std or Phe buffer was used as mobile phase using a flow rate of 2.5 mL·min−1, and UV‐detected (280 nm) peaks containing the target protein  were collected. The concentration of the collected protein was determined by UV‐Vis absorption at 280 nm utilizing theoretical extinction coefficients obtained from Expasy Evaluation of molecular weights and purity were conducted by SDS/PAGE. Proteins were resolved on Mini‐PROTEAN TGX gels (7.5%) from Bio‐Rad  run at 100 V for 75 min with a protein standard from Bio‐Rad (no. 161‐0304). Gels were stained with Coomassie Blue to visualize the proteins.m, SYPRO orange at a concentration of 2× , and ligand (diluted in purification buffer) in a concentration range of 0.1 μm to 10 mm. Each plate contained control wells with the purification buffer with and without protein and ligand. Scans were carried out using a scan rate of 1 °C·min−1, going from 20 °C to 95 °C. The thermograms were baseline corrected with mxpro qpcr Software  and analyzed for transition temperatures with graphpad prism 6  utilizing a Boltzmann sigmoid fit:a denotes the slope of the unfolding curve at TmThe unfolding of the TPH2 variants was recorded with an Agilent Technologies Stratagene MX3005 P RT‐PCR machine . The ligand screen was performed with a total volume of 25 μL in 96‐well plates (polypropylene plates from Agilent Technologies). Each well was composed of protein at a concentration of 1 μl‐Trp, l‐Phe, or d‐Phe and allowed to equilibrate for 20 min. Samples were injected with a 500 μL loop and analyzed at a flow rate of 0.5 mL·min−1. Calibration curve of molecular weights was obtained from GE Healthcare Determination of the oligomeric state of the TPH2 variants was performed on a Superdex 200 10/300 GL column. Prior to analysis, the column was equilibrated with two column volumes of the buffer of investigation, and the samples were spiked with concentrated stock solutions of The equilibrium constant of the dissociation of a dimer into monomers was based on the scheme belowKd, is defined bytotal, can be expressed in terms of molar monomer equivalents280,dimer = 2·ε280,monomer and area under the size exclusion curve (SEC) curve (AUC) ∝ [M], the molecular concentration of monomer is given byM and AUCD are the area under the curve of the peaks representing the monomer and dimer, respectively, obtained by deconvolution of the chromatograms from the SEC experiments. Substituting [M] in equation D/(AUCM + AUCD) is the mole fraction of dimer, fD. Substituting [M] and [D] from equation total must be divided by the dilution factor, DF, introduced during gel filtrationfD, as a function of [M]total can be expressed asThe equilibrium dissociation constant, fD, yieldsSolving equation hTPH2 variants were thawed under running water, filtered, and the concentrations were determined by UV‐Vis absorption at 280 nm. The hTPH2 samples were diluted to a protein concentration of 5 μm in the buffer in which it was purified. TPH2 activity was assayed in a reaction mixture  – 2500 μL total volume) containing 50 mm HEPES/NH4OH, 200 mm (NH4)2SO4, pH 7.0, 0.025 g·L−1 catalase, 25 μm (NH4)2Fe(II)(SO4)2.6H2O, 7 mm dithiothreitol (DTT), 60 μm l‐Trp, and 300 μm BH4 with stirring at 15 °C −1) of the monitored fluorescence and were quantified using varian spectrophotometer software (Agilent Technologies). Quantification occurred through linear regression on a manually placed interval of minimum 0.04 min of the initial curve. Denaturation was performed by heating the protein in aliquots of 1300 μL to 30 °C in a water bath. The denaturation was stopped by cooling the protein solution in ice water.The activity measurements were performed using a Varian Cary Eclipse Fluorescence Spectrophotometer. For activity measurements, the l‐Phe containing buffer were used for sample measurements, and the same buffer as used in the purification of the hTPH2 variants were used as references. Prior to sample loading, the protein sample and the reference buffer were degassed at approximately 10 °C using a MicroCal  USB ThermoVac. Three hundred microliters of protein sample and reference buffer were loaded in the respective cells. A constant pressure of 3 atm was applied, and scans were carried out using a scan rate of 1 °C·min−1, going from 5 °C to 70 °C. Deconvolution of the thermograms was performed with Peakfit v4.12 using Gaussian peak functions.Differential scanning calorimetry (DSC) measurements were carried out on a TA Instruments  6300 Nano DSC. Solutions of protein in KDT: Wrote the paper, performed experiments, and analyzed the data. HEMC and GHP: Designed experiments, contributed to data analysis and in writing the paper. NH and JB: Performed preliminary experiments. PH: Contributed to writing the paper.Appendix S1. Sequences of hTPH2 protein variants.Table S1. Primers for cloning the recombinant truncated hTPH2 protein variants.Fig. S1. Alignment of the regulatory domains of hTPH2 and rnPAH.Fig. S2. Representative example of raw data from differential scanning fluorimetry.Fig. S3. SDS‐PAGE results of collected peaks from purification of rchTPH2 using Phe buffer.Fig. S4. SDS‐PAGE results of collected peaks from purification of NΔ47‐rchTPH2 using Phe buffer.Fig. S5. SDS‐PAGE results of collected peaks from purification of chTPH2 using Phe buffer.Click here for additional data file.


The designed unique architecture allows efficient use of pseudo‐capacitive Mn2O3 nanomaterials for charge storage with facilitated transport for both ions and electrons, rendering them high specific capacitance, good rate capability, and remarkable cycling performance. Among them, the manganese oxide is one of the most promising candidates due to its high specific capacitance, superior environmental compatibility and cost effectiveness.2 (Mn[IV]),[[qv: 5a]] Mn2O3 (Mn[III])[[qv: 5b]] or Mn3O4 (Mn[II/III]).[[qv: 5c]] Though manganese oxides have demonstrated huge potentials as supercapacitor electrodes, they still suffer from some drawbacks:[[qv: 2e]],[[qv: 2f]],[[qv: 4a]],Supercapacitors are intermediate devices between conventional batteries and dielectric capacitors that can be used in various electric systems.An emerging solution to overcome the above drawbacks is to utilize the unique structural characteristics of the hollow micro‐/nano‐structured materials.3O4,[[qv: 9a]] Fe2O3,[[qv: 9b]] SnO2[[qv: 10c]] and MFe2O4[[qv: 10a]] ) with elaborate control over the number of internal shells, shell spacing, and exterior shell structure. To date, however, the multi‐shelled hollow structured MnOx materials have not been used for supercapacitor applications mainly due to the difficulties in using the template methods to achieve the control of composition, crystalline phase and microstructure  of MnOx.[[qv: 2a]],],[[qv: 4a]]Recently, a general method utilizing the carbonaceous microspheres (CMSs) as sacrificial templates has been developed by our group for controllable synthesis of hierarchical multi‐shelled hollow microspheres with adjustable nanoparticle subunits and shell structures.[[qv: 8a]],2O3 hollow microspheres with controlled number of shells, shell thickness and porosity as the high performance supercapacitor electrode materials. The as‐synthesized triple‐shelled Mn2O3 hollow microspheres with thin porous shells show an extremely high specific capacitance up to 1651 F g–1 at a current density of 0.5 A g–1, and remarkable cycling stability with 92% retention after 2000 consecutive cycles. Besides, the rate capability is also impressive, showing a specific capacitance as high as 1422 F g–1 at a high current density of 10 A g–1. To our knowledge, the performance of Mn2O3‐based electrode materials for supercapacitor reported here is a new record to date.Herein, we report an effective approach to synthesize the multi‐shelled MnAs previously reported, the control over the crystalline phase and microstructure  can be achieved by tuning the metal ions adsorption amount and the depth embedded into the carbonaceous microsphere templates, as well as tuning the annealing conditions.[[qv: 8a]],Figure2 precursor solution at 3.35, 4.43 and 6.43, respectively, the single‐, double‐ and triple‐shelled hollow microspheres can be obtained in a high yield after the removal of the templates by heat treatment . It is well known that a decrease in pH leads to a decrease in the surface negative charge of CMSs, hence a decreased Mn cations adsorption onto the CMSs template.[[qv: 10c]] As shown in Figure S1, Supporting Information, the zeta potential of carbonaceous microsphere templates in aqueous solution is about –44.5 mV at the pH value of 6.78. With the solution pH values of 3.52, 4.80 and 6.23, the corresponding zeta potentials of CMSs are found to be –20.4, –29.9, and –42.0 mV, respectively, meaning a less electrostatic attraction between the negatively charged CMSs and the Mn cations. Moreover, with the addition of HCl aqueous solution (pH regulator), the coordination group of Mn2+ ions are changed from [Mn(H2O)6]2+ to [Mn(H2O)6‐xClx]2‐x , further decreasing the electrostatic attraction.2O3 hollow microspheres with certain number of shells can be obtained after heat treatment  instead of Mn2O3 hollow microspheres .In this communication, we report a modified control synthesis approach for fabrication of uniform single‐, double‐, triple‐shelled hollow microspheres with thin porous shells by precisely controlling the pH values of the precursor solutions with an appropriate concentration. As shown in 2O3 hollow microspheres can be obtained  into the aqueous infusion solution . According to the Arrhenius equation,[[qv: 9a]],2O3 hollow microspheres and nanoparticles were investigated by X‐ray diffraction (XRD). As shown in Figure 2O3 phase (JCPDS card no. 41–1442), with no additional peaks are detected.x that exist in the resulting products, such as MnO2, probably due to there not being enough driving force to step over the energy barrier to a much higher valence state. Further compositional information can be obtained by Raman spectra , EDX and corresponding elemental mapping . The average grain sizes of Mn2O3 nanoparticles and hollow microspheres with single‐, double‐, triple‐, and quadruple‐shells calculated by Scherrer Equation are 19.54, 30.35, 28.43, 27.12 and 29.87 nm respectively, consistent with the TEM (transmission electron microscopy) and SEM (scanning electron microscopy) observations shown in Figure 2O3 phase with high crystallinity, which is also proved by the high‐resolution TEM (HRTEM) image . Taking triple‐shelled Mn2O3 hollow microspheres as a representative example, two characteristic broad peaks encompassing the Mn (III) 2p3/2 at 641.8 eV and 2p1/2 at 653.3 eV are observed, with a spin‐energy separation of 11.5 eV, which is consistent with the reported data of Mn (III) 2p3/2 and Mn (III) 2p1/2 in pure Mn2O3 phase.The crystalline structures of as‐prepared Mn2O3 hollow microspheres as electrode materials for supercapacitors were evaluated by a three‐electrode electrochemical system.Figure2O3 hollow microspheres electrodes obtained from 6 m KOH electrolyte under scan rates ranged from 5 mV s–1 to 200 mV s–1 within the potential window of 0.0–0.45 V ). For comparison, the electrodes assembled with single‐, double‐, quadruple‐shelled Mn2O3 hollow microspheres and Mn2O3 NPs are also evaluated . It can be found that the CV curves of Mn2O3 are distorted from mirror image symmetry and a pair of redox peaks can be clearly observed, which can be ascribed to the redox reactions between the oxidation states of Mn(III) and Mn(IV). Therefore it can be known the Faradic capacitance plays the main role in the charge storage of Mn2O3 hollow structures. Moreover, all these CV curves preserve the obvious redox peaks of Mn2O3 even at a high scan rate of 200 mV s–1, indicating an ideal capacitive behavior and desirable fast charge/discharge property for power devices.The performance of the synthesized multi‐shelled MnI–V responses and small equivalent series resistance, ideal for capacitive device applications. The specific capacitance of the triple‐shelled Mn2O3 hollow microspheres electrode was found to decrease with increasing discharge current densities. It can be seen from Figure S9, Supporting Information, that the specific capacitance of triple‐shelled Mn2O3 hollow microspheres was 1651 F g–1 at a discharge current density of 0.5 A g–1 and still kept as high as 1422 F g–1 at a high discharge current density of 10 A g–1, indicating a superior rate performance. Remarkably, compared with other Mn2O3‐based supercapacitor electrode reported up to date, the specific capacitance of the triple‐shelled Mn2O3 hollow spheres is a new record .To further evaluate the cell performance, the galvanostatic charge/discharge curves at various current densities were measured with a constant voltage window of 0.0–0.45 V.2O3 hollow microspheres can be well controlled to optimize the capacitive performance. Figure2O3 NPs and hollow microspheres with different shell numbers (1–4) to explore the effect of shell numbers on the resultant capacitances. It can be seen that the CV integration areas per weight of triple‐shelled Mn2O3 hollow microspheres, which indicates the charge stored in the supercapacitor, are much larger than that of Mn2O3 NPs, single‐, double‐ and quadruple‐shelled hollow microspheres. It seems that the maximum capacitance could only be obtained from a moderate shell number of Mn2O3 hollow microspheres with the porous thin shell. This is easy to be understood that as the maximal active sites and appropriate pore volume could only be obtained from the proper shell structures to facilitate the ions penetration and the process of redox reactions. The specific capacitance of the electrodes with different scan rates is given in Figure 2O3 hollow microspheres electrodes were much higher than that of the Mn2O3 NPs, and among them triple‐shelled Mn2O3 hollow microspheres achieved a maximum capacitance, the crest value of specific capacitance of 1640 F g–1 at the scan rate of 5 mV s–1 can be achieved with the triple‐shelled Mn2O3 hollow microspheres. The same trend can also be observed by comparison of galvanostatic charge/discharge curves of various samples at the same charge/discharge current density . The Brunauer–Emmett–Teller (BET) data  show that the triple‐shelled Mn2O3 possesses the highest specific surface area of 36.55 m2 g–1 and also the largest pore volume of ≈0.518 cm3 g–1 among other multi‐shelled hollow microspheres investigated . It can be deduced that a higher specific surface area and larger pore volume could guarantee an abundance of the active sites and the easy electrolyte penetration into the materials, resulting in an improved capacitive performance. It should be mentioned that the trend of the specific capacitance changes is consistent with the changes of specific surface areas and the pore volumes of multi‐shelled Mn2O3 hollow microspheres, supporting our standpoints. However, when comes to Mn2O3 NPs, it shows a much higher specific surface area and larger pore volume, yet a lower specific capacitance. It could be explained from two aspects: first, although Mn2O3 NPs possess a larger surface area, the actual accessible reactive surface area of hollow microspheres is much larger due to the special hierarchical structure and well‐preserved interface dispersibility in the electrode layers, while the Mn2O3 NPs could lose large portion of their reactive sites due to serious accumulation and crosslink. Additionally, the unique three‐dimensional hierarchical porous multi‐shelled hollow structure can ensure the ions from both outer and inner shell to be charged and transferred, resulting in superior capacitive and rate performance.Since the charge storage process of a supercapacitor is a surface reaction between the electrolyte ions and electrode materials,[[qv: 2b]] a detailed surface structure investigation is therefore needed to understand the mechanism behind the discrepant performance. From the TEM and SEM images we can clearly see that the shell thickness of multi‐shelled hollow microspheres is different and triple‐shelled Mnx often suffers from a limited long‐term stability because of the partial dissolution in the electrolyte during cycling, which restricts the commercial application of such low‐cost supercapacitor electrode materials.x hollow microspheres is able to noticeably meliorate the cycling performance. The cycling life test of Mn2O3 NPs and hollow microspheres were examined between 0 and 0.45 V in 6 m KOH aqueous electrolyte by consecutive charge/discharge cycles at a current density of 1 A g–1 , while Mn2O3 NPs could only retain a capacitance of 376.5 F g–1 with a 35.3% loss of its initial capacitance (581.7 F g–1). The better stability of the multi‐shelled Mn2O3 hollow microspheres can be ascribed to the superior shell structure with the shells supporting for each other, especially the exterior shell protecting the interior shells, which can undertake some mechanical deformation and swell in the redox processes and suppress the dissolution of Mn2O3 in electrolyte.Long cycling life is also an important requirement for supercapacitors. In this regards, MnO2O3 hollow microspheres as high‐performance supercapacitor electrodes and the effect of shell structures on the capacitive performance are carefully studied. The unique architecture of these designed nanostructures allows efficient use of pseudo‐capacitive Mn2O3 nanomaterials for charge storage due to the facilitated transport of both electrolyte ions and electrons, thus rendering them with a record high specific capacitance of 1651 F g–1, good rate capability and remarkable cycling performance. We believe that such earth‐abundant, low‐cost, environment‐friendly, and high‐performance electrode materials achieved through a scalable and controllable process, can offer great promise in large‐scale energy storage device applications.In summary, we have developed a scalable approach to fabricate multi‐shelled MnAs a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


This suggests that zinc can react with high-valence sulfur to generate S2− at digestion temperature, which is consistent with our thermodynamic calculation results. In this study, as zinc dosage increases, sulfur digestion rate decreases while sulfur content in red mud markedly increases when zinc dosage was below 4%; the digestion rates of sulfur and sulfur content in red mud remains stable when zinc dosage was above 4%; the alumina digestion rate, conversely, increased slightly throughout the experiment. This suggests that high-valence sulfur in sodium aluminate solution can be converted to S2− and then enter red mud to be removed completely by adding zinc during the digestion process.This paper proposes a novel approach to sulfur removal by adding zinc during the digestion process. The effects of zinc dosage on the concentrations of different valence sulfur in sodium aluminate solution were investigated at length to find that high-valence sulfur (S China possesses over 800 million tons of diasporic high-sulfur bauxite, mainly in Henan, Guizhou, and Chongqing provinces2−, then the S2− is gradually oxidized into various forms of high-valence.During the Bayer process, the sulfur in high-sulfur bauxite first enters the solution in the form of S2O32−, SO32−, SO42−). The most negative effects of sulfur’s presence in the Bayer process are as follows3: 1) Na2S can react with Fe2O3, and Na2S2 can react with Fe(OH)2 to form Na2[FeS2(OH)2]·2H2O, which is much more soluble in sodium aluminate liquor; this increases the iron content in alumina products. 2) Na2S2O3 can react with Fe to form Na2S, Na2SO3, and Fe(OH)2, which destroy the steel oxide film and accelerate the corrosion of equipment. 3) The sulfur in the Bayer process is eventually converted to Na2SO4, which increases alkali consumption. 4) Once the content of Na2SO4 in spent liquor reaches a certain level, it can precipitate in the form of Na2CO3·2Na2SO4; the double salt can then scale the evaporator and interior digester, which impact industrial production.Sulfur . These methods have their own advantages and disadvantages, they have not been widely applied in industry. Beneficiation desulfurization prevents most of the sulfur in bauxite from contaminating the alumina production process, but sulfur in the concentrate remains problematic, requiring further desulfurization throughout the process. Wet oxidation desulfurization does not introduce any new impurities into the liquor, and the necessary materials are cheap and readily available, but some amount of hydrogen is produced during wet oxidation, so there are safety risks inherent to the process, further, discharge and causticization increase the production cost. Desulfurization efficiency of desulfurization by precipitators is high, but only a form of sulfur was removed.To date, relatively few researchers have investigated sulfur removal in the alumina production process because high-sulfur bauxite is seldom used in alumina production outside of China, but there are many studies of sulfur removal in other fields2− from oxidizing into S2O32−, SO32−, and SO42− in the sodium aluminate solution, which forces S2− into red mud to decrease the concentration of S2−, S2O32−, SO32−, SO42−, and total sulfur in the liquor. As opposed to other methods of sulfur removal, this method does not need discharge or causticization processes to remove the sulfate, and all forms of sulfur in liquor can be removed. As discussed below, removal of sulfur via the zinc addition in liquor was investigated to provide a theoretical and technical basis for the effective utilization of high-sulfur bauxite.Large swaths of China’s high-sulfur bauxite reserves are unable to be developed or utilized due to the lack of appropriate deep desulfurization method for alumina production. This paper proposes a new method of sulfur removal in which zinc is supplied to prevent Sk (molar ratio of Na2Ok to Al2O3) of 1.42. The effects of zinc dosage on the concentrations of different valence sulfur in sodium aluminate solution are shown in Fig. Zinc dosage was varied between 0 and 5% through a series of digestion experiments at temperature of 533 K (260 °C), digestion time of 60 min, lime dosage of 13%, digestion liquor α2O32−, SO32−, SO42−) in the sodium aluminate solution decreased as zinc dosage increased, while the calculated values of the S2− concentration in the solution increased notably with the increase of zinc dosage. This suggests that zinc can react with high-valence sulfur to generate low-valence sulfur at the digestion temperature, which is consistent with our thermodynamic calculations  was calculated using Factsage 7.0 software, the results are shown in Fig. The standard Gibbs free energy of the reactions of zinc with high-valence sulfur , the ΔGθ of the reaction of zinc with SO42− was negative value when the temperature above 448 K (175 °C). This suggests that zinc can react with high-valence sulfur to generate S2− at digestion temperature, because digestion temperature of high-sulfur bauxite is higher than 513 K (240 °C) in industrial production.As can be seen from Fig. Gθ value, the more favorable the reaction is, so S2O32− is easiest to be reduced by zinc, then SO32−, SO42− is most difficult to be reduced at temperature of 533 K (260 °C), this explains why S2O32− in liquor were removed completely, the SO32− and SO42− were removed nearly completely when zinc dosage was 5%  were tested under the same other conditions described above. The Effects of zinc dosage on the digestion rates of alumina and sulfur are shown in Fig. 22−, the solubility product constant of ZnS (2 × 10−24) is very low, so it is difficult to decompose in sodium aluminate solution, the S2− can react with Zn2+ easily to generate ZnS34, the more S2− in liquor, the easier ZnO22− produce to Zn2+. As a reductant, the zinc reacted with high-valence sulfur to generate S2−, and then the S2− reacted with Zn2+ to generate ZnS went into red mud, so the sulfur in liquor was removed. The reactions are as follows:It can be seen from Fig. The sulfur content increased in red mud as zinc dosage increased  were tested under the same other conditions described above. The Effects of zinc dosage on the concentrations of iron in sodium aluminate solution are shown in Fig. 2O3 decreased substantially as zinc dosage increased; the concentration of Fe2O3 remains stable when zinc dosage was above 4%. So iron in sodium aluminate solution was also removed, when sulfur in liquor was removed by adding zinc during the digestion process.It can be seen from Fig. Again, six different zinc dosages (between 0% and 5%) were tested under the same other conditions described above. The effects of zinc dosage on absorbance of liquor are shown in Fig. It can be seen from Fig. It can be seen visually from Fig. 35. So, it can be seen from Fig. The higher the sulfur concentration in sodium aluminate solution, the deeper the colour of solution2O32−, SO32−, SO42−) to generate low-valence sulfur (S2−), zinc in sodium aluminate solution mostly in the form of ZnO22−, the S2− promotes ZnO22− to produce Zn2+, and then the S2− reacted with Zn2+ to generate ZnS went into red mud. As a reductant and precipitator, the zinc can remove sulfur completely in sodium aluminate solution during the digestion process.Based the above result and discussion, we propose the following mechanism of sulfur removal  in the filtrate were chemically analyzed36. The dried red mud was sampled and observed with an X-ray fluorescence analyzer and carbon-sulfur analyzer.The mineral sample and evaporation spent liquor according to a certain proportion were placed in a 100 ml stee1 bomb which was sealed in the experiment. The digester was heated to 533 K (260 °C) and held for 5 min, then the steel bomb was placed inside and the digestion process was run 60 min. The sodium aluminate solution was then filtered and the AlA) was calculated as follows:The digestion rate of alumina mud Mass ratio of alumina to silica in mud(A/S)S) was calculated as follows:The digestion rate of sulfur (ηore Mass percentage of sulfur in raw ore (%)Smud Mass percentage of sulfur in mud (%)Sore Mass percentage of iron in raw ore (%)Fmud Mass percentage of iron in mud (%)F