Nucleolytic ribozymes utilize general acid‐base catalysis to perform phosphodiester cleavage. In most ribozyme classes, a conserved active site guanosine is positioned to act as general base, thereby activating the 2′‐OH group to attack the scissile phosphate . Here, we present an atomic mutagenesis study for the pistol ribozyme class. Strikingly, “general base knockout” by replacement of the guanine N1 atom by carbon results in only 2.7‐fold decreased rate. Therefore, the common view that γ‐catalysis critically depends on the N1 moiety becomes challenged. For pistol ribozymes we found that γ‐catalysis is subordinate in overall catalysis, made up by two other catalytic factors (α and δ). Our approach allows scaling of the different catalytic contributions  with unprecedented precision and paves the way for a thorough mechanistic understanding of nucleolytic ribozymes with active site guanines. General base knockout in nucleolytic ribozymes using deazaguanosine‐modified mutants paves the way to scale the relative contributions of the four main strategies  applied in ribozyme general acid base catalysis. Ka values far from neutrality and so are not optimal for general acid‐base catalysis.Ka values[+, C+, and G−, respectively, although the corresponding free nucleotides are neutral in aqueous solution at pH 7.Small self‐cleaving ribozymes are crucial players in the life cycle of cellular RNAs, for example, they control replication of satellite and pathogenic RNAs.2, 3 FThe catalytic contributions to phosphodiester cleavage have been categorized into for major strategies Figure A: 1) In‐2+ ion, innersphere‐coordinated to the N7 of G33, stands out through participation as general acid.Pistol ribozymes Figure A,B repreIn contrast, the role of the general base is less clear. Although all available crystal structures suggest an active site guanosine (G40) to embark on this role,32, 331G), 3‐deazaguanosine (c3G), and xanthosine (X) needed for RNA solid‐phase synthesis. Since only for c3G containing RNA synthetic procedures were published,[1G and X modified RNA.[To achieve our goals, synthetic hurdles had to be taken first; those concerned access to the phosphoramidites of 1‐deazaguanosine  substrate. This is consistent with the crystal structure, showing a distance of about 48 Å distance between 5′‐ and 3′‐termini of the substrate strand in stretched conformation .Figure e Figure B. This r2+ addition  compared to the wildtype G40 ribozyme . Clearly, c1G does not offer these functionalities and therefore, a much higher impact on activity was expected upon N1‐by‐C1 replacement. In this context, we note that the pKa of c1G is 9.1,Ka value is comparable to the pKa of G and it is tempting to speculate that the deprotonated O6 in c1G may take over the role of a deprotonated N1 in G although it is dislocated compared to the N1 position. We therefore tested a pistol mutant with 1‐deazainosine  is 12.3[Ka 9.2). This certainly affects H‐bonding properties and the propensity for activation of the 2′‐O nucleophile. Indeed, the c3G40 mutation had a larger impact on cleavage compared to c1G40, reflected in a 7.2‐fold decrease in the observed rate  and H2N(C4) of C41 is impaired,[Ka.To further shed light on the role of G40 in pistol ribozyme cleavage, we synthesized the ce Figure B. Nevert2 (for G) to H‐acceptor properties through the C2 carbonyl group (for X). Given the generally accepted opinion that in many nucleolytic ribozymes the C2‐NH2 group of the guanine general base is crucial for stabilizing the pentavalent phosphorane transition state by H‐bond interactions with one of the non‐bonding oxygens,[2 and the non‐bridging oxygen observed in the crystal structure of the transition state mimic ,[Next, we investigated the impact of xanthosine in position 40 on pistol activity. This mutation changes the H‐donor properties of the original exocyclic NH oxygens,43, 44c Figure C. We furKa values of the titratable functional groups in the catalytic pocket.To shed further light on the specific impact of G40 as general base in pistol ribozyme cleavage we recorded activity‐pH profiles of the wildtype and the diverse mutant ribozymes by additional experiments at pH values ranging from 5.0 to 9.5 Ka of 6.94 assigned to the influencer, 7.99 for the general base, a ΔpKcoop of 1.94 and 8.9 for the general acid (Table Ka for the influencer remained slightly higher than observed in the data points (rounding point ≈6.5) whereas the pKa values for the general base and acid are predicted marginally lower compared to the actual rounding points of the measured profile (>8 for the base and >9 for the acid). Further, the coupling ΔpKcoop between the influencer and the general base of 1.94 is predicted higher than observed in the first plateau, that extends only over 1 pH unit from 6.5 to 7.5. Nevertheless, we clearly see that the pKa of the general base is significantly shifted to neutrality by ΔpKcoop units in the presence of an active protonated influencer species. Simultaneously, the activity of the influencer is enhanced by an upward shift of ΔpKcoop units towards the neutral range. Even in the absence of the functionally active influencer above pH 8, pKa shifts for the general base of about 1.2 pKa units and about 2 pKa units for the general acid are observed compared to the unshifted values  lose the wavy shape of their rate‐pH profiles. They seem to lack cooperative interactions with an influencer, resulting in pKa values in the same range as modelled for the wildtype in absence of an influencer. Finally, we note that also for the pistol ribozyme from Paenibacillus polymyxa, a “wavy” rate‐pH profile was observed although this feature was not further discussed in that study.Interestingly, the G40 variants tested here . Therefore, even if G42 would be able to take over the role of G40, its contribution to catalysis is minor. We further note that G42 bridges G40 and A32 via an H‐bond network  and thereby generates a cleft to accommodate the scissile phosphate of the splayed‐apart G53–U54 nucleosides .[1G mutation retains these interactions while other modifications, such as 2‐aminopurine (Ap), do not. The significant rate reduction previously observed for G42Ap mutated pistol ribozyme may arise from this structural impairment.In a recent computational study, a possible catalytic role of an alternative conserved guanosine in the binding pocket, namely G42, has been postulated.re S11A).30, 347 to 108).Nucleolytic ribozymes accelerate the cleavage reaction by more than a million‐fold 1007 to 1082+ was mutated to C7 (c7G33), the cleavage rate was reduced by an order of at least 103.[2+ as general acid provides a significantly higher contribution compared to γ‐catalysis of the general base G40. The question still is where the remaining factor of about 103 to 104 to achieve an overall acceleration of 107 to 108 originates from. This acceleration apparently stems from α‐ and β‐catalysis. For pistol, the contribution through transition state stabilization  is also small as deduced from comparatively small rate reductions by c3G (6‐fold), xanthosine (53‐fold), and 1‐deazainosine  to P−O5′ (of nucleotide N+1) according to Ref. τ close to 180°. Additionally, our FRET folding assay indicates that the extended conformation of the substrate is already adopted prior attack of the nucleophile at the scissile phosphate. Notably, given the prominent role of substrate positioning in the ribosome and of cofactor positioning in recently found methyltransferase ribozymes,[The present study provides scaling of catalytic contributions for the pistol ribozyme class with unprecedented precision. Although it is tempting to assume that in general γ‐ and δ‐catalysis—jointly constituting concerted general acid‐base catalysis—make the largest contribution to the catalytic rate enhancement,east 103.33, 341–102) to the estimated overall rate enhancement in pistol ribozymes. Even G40rS (lacking the nucleobase completely) and the double mutant G40c1G–G42c1G have comparable residual cleavage activities  .Finally, we note that the strict phylogenetic conservation of pistol G40 still makes sense from the point of evolution, aiming at utmost activity attainable from the four naturally available nucleotides. Pistol mutants G40A, G40C, and G40U have 34‐fold, 94‐fold, and 152‐fold less cleavage activity . This is consistent with the small contribution of β‐ and γ‐catalysis of nucleotide‐40 ,To date, only one other ribozyme class has been scaled for catalytic contributions with high precision. This is twister where “general base knockout” using c1G. This feature of a “general base knockout” has been neglected in earlier mechanistic studies of nucleolytic ribozymes. Our study paves the way for scaling catalytic factors in other ribozymes as well, and therefore, contributes to a thorough mechanistic understanding of ribozyme catalysis.In summary, our study demonstrates a reliable approach for scaling the different contributions of catalysis  in individual self‐cleaving ribozymes. This becomes possible by directly addressing the key functionality N1 of the general base guanine through atomic mutagenesis with cThe authors declare no conflict of interest.1As 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.Supporting InformationClick here for additional data file.


But this property is still limited by the low electronic conductivity or insufficient wiring manner. Herein, a strategy of triple conductive wiring through electron doping, chelation coating, and electrochemical conversion inside the microsized porous spheres consisting of dendrite‐like T–Nb2O5 primary particles is proposed to achieve the fast‐charging and durable anodes for LIBs. The penetrative implanting of conformal carbon coating (derivative from polydopamine chelate) and NbO domains (induced by excess discharging) reinforces the global supply of electronically conductive wires, apart from those from Co/Mn heteroatom or O vacancy doping. The polydopamine etching on T–Nb2O5 spheres promotes their evolution into fluffy morphology with better electrolyte infiltration. The synergic electron and ion wiring at different scales endow the modified T–Nb2O5 anode with ultralong cycling life (143 mAh g−1 at 1 A g−1 after 8500 cycles) and high‐rate performance (144.1 mAh g−1 at 10.0 A g−1). The permeation of multiple electron wires also enables a high mass loading of T–Nb2O5 (4.5 mg cm−2) with a high areal capacity of 0.668 mAh cm−2 even after 150 cycles.High‐rate anode material is the kernel of developing fast‐charging lithium ion batteries (LIBs). T–Nb 2O5 primary particles is proposed to achieve the fast‐charging and durable anodes. The synergic electron and ion wiring at different scales endows the modified T–Nb2O5 anode with ultralong cycling life and high‐rate performance.A strategy of triple conductive wiring through electron doping, chelation coating, and electrochemical conversion inside the microsized porous spheres consisting of dendrite‐like T–Nb However, generally the practice of fast‐charging either sacrifices the energy density and mileage range by compromising the cell‐design metrics or risks the lifespan and safety of batteries otherwise. Current research efforts are thus attracted to the exploitation of high‐power electrode materials, which directly tackles the technical bottleneck.Although the gradual proliferation of electric vehicles (EVs) is an unstoppable trend into our future, it remains a puzzle how to fill the remaining gap in performance between EV batteries and combustion engines. In particular, the criterion of US Advanced Battery Consumption (USABC) for high performance commercial batteries stipulates 80% state of charge achieved within 15 minutes during a regular recharge, which is still a challenge.1, 2 Though graphite displays the relatively high specific capacity (372 mAh g−1) and highly reversible (de)intercalation performance, it is widely deemed to be the culprit that impedes the achievement of fast‐charging due to its low intercalation potential (≈0.1 V vs Li+/Li). As a promising alternative anode material, orthorhombic niobium pentoxide (T–Nb2O5) is well known for its intercalation‐based pseudocapacitive effect, enabling its high‐rate charge storage. Inside the distinct layered structure of T–Nb2O5 with “room and pillar” NbO6/NbO7 units, the covalent bonds (Nb–O–Nb) instead of weak van der Waals forces  connect the adjacent niobium atom layers. This unique structure enables the formation of robust Li diffusion channels and prevents structural exfoliation as observed in graphite under ultrafast Li+ (de)intercalation. Furthermore, Li deposition is another well‐known limitation for graphite anode during high‐rate charging due to its low intercalation potential. This issue can also be circumvented in T–Nb2O5 considering its intercalation potential range (1.0–3.0 V). Moreover, micrometer‐scale Nb2O5 particles also display higher tap density than commercial graphite, Li4Ti5O12 and TiO2, promising an opportunity to further improve the cells’ volumetric energy density. It deserves to highlight that no requirement on complicate nanostructure design on Nb2O5 can mitigate the fracture of fragile solid electrolyte interface and avoid the unconscionable Li depletion. These prominent and luciferous prospects of T–Nb2O5 have attracted more and more researchers into the field, but it is also plagued by several intractable disadvantages  when compared with commercial graphite anode.Graphite anode is the foundation for the commercialization of lithium‐ion batteries (LIBs) in 1991.4 Thou2O5 is identified as an electrical insulator (with a band gap of 3.2–4.0 eV) and it can be converted to n‐type semiconductor with the implantation of trace amount of oxygen vacancies. The intrinsically inferior electric conductivity of Nb2O5 (≈3 × 10–6 S cm−1) would trigger the particle pulverization and rapid capacity decay during the iterative charge/discharge process. A variety of methods have been done to tackle this problem. For instance, a nitrogen‐doped amorphous carbon layer has been demonstrated to homogenize electron transport in monoclinic Nb2O5 (H–Nb2O5) so as to suppress inhomogeneous phase change, finally achieving excellent rate performance (≈120 mAh g−1 at 16.0 A g−1). Besides, Cu2Nb34O87, a compound based on Cu2O doped Nb2O5, exhibits enhanced mass and charge transport (3.5 × 10–13 cm2 s−1 for Li+ and 2.1 × 10–5 S cm−1 for electron). Analogously, WO3–Nb2O5 solid solution (Nb14W3O44) was found to facilitate the construction of interconnected 3D Li+ migration (10–10–10–12 cm2 s−1) tunnels in microsized particles. Synergistic engineering of electron and ion transport pathways appears to be effective in promoting the rapid and reversible Li+ (de)intercalation in niobium‐based oxides. Exterior conductive carbon wiring and interior defect‐rich structuring synergistically endow Nb2O5 particles with ultrafast and stable Li storage performance. By far, the majority of studies focus on the modification of closely packed structures, such as orthorhombic phase, high temperature phase and ReO3‐type Wadsley–Roth phases, constructed with corner‐shared or edge‐shared octahedrons and tetrahedrons. Compared with graphite, the relatively dense and rigid structure of Nb2O5 is less conducive to the rapid Li migration inside the layered structure because of the higher steric hinderance and higher migration barrier. Enlarging the interlayer spacing is beneficial to the charge storage kinetics in intercalation‐type materials. The electrochemical kinetics of Nb2O5 based anode still remains to be optimized in such aspects as the homogenization of electron transport and construction of stable Li+ channels.In general, stoichiometric Nbacancies.9 The 2O5 to achieve a fast‐charging and durable anode for LIBs. This T–Nb2O5–based anode is firstly doped by Co or Mn heteroatoms to induce the creation of oxygen vacancies to improve its electron conductivity. Afterward, a polydopamine (PDA) layer chelates with and etches the T–Nb2O5 surface, and it is then pyrolyzed into the exterior carbon coating to further increase the electron conductivity of doped T–Nb2O5. The PDA etching also triggers the evolution of irregular spheres of T–Nb2O5 into the well‐defined porous spheres consisting of dendrite‐like primary particles. This fluffy morphology can promote the electrolyte infiltration and sufficient Li‐ion intercalation. The excessive injection of Li‐ions into the lattices of this anode by discharging to a lower voltage allows the in‐situ wiring of conductive NbO domains (with a reported electron conductivity of 4.8 × 104 S cm−1 at 300 K) inside the electrochemically cycled T–Nb2O5, as a result of the phase segregation between NbO and Li‐stuffed LixNb2O5. This new discovery and successful detection of NbO domains explain the much better capacity retention than the case without prior deliberate over‐discharge. The above triple wiring design reinforces the global supply of electron and Li‐ion flows, endowing the modified T–Nb2O5 anode with ultralong cycling life and high‐rate performance, as well as the potential of high mass loading and high areal capacity. The Li storage mechanism related to the formation and evolution of NbO wiring network is discussed based on in situ structure characterization and electrochemical kinetics analysis in detail.In this paper, we propose a comprehensive strategy of triple conductive wiring through electron doping, chelation coating and electrochemical conversion inside T–Nb22O5 is usually limited to 1–3 V when it is adopted as an anode material in LIB. With the ongoing Li intercalation in the discharge process, the Nb2O5 particles can become electronically conductive. Here we deliberately broaden the electrochemical window to 0.4–3 V to induce the formation of electronic conduction network inside the micrometer‐scale T–Nb2O5 particles. The lower cut‐off potential (<0.4 V) is not desirable in view of the risk of Li metal deposition especially under high current density. Figure xNb2O5 to the amorphized hybrid of NbO, Li2O, and LixNbyOz. In the voltage range of 1–3 V, the reversible Li storage in T–Nb2O5 mainly proceeds through a solid solution reaction behavior  and PDA‐derived carbon coated Co–NbO (denoted as PDA–Co–NbO). During the solvothermal reaction in isopropanol, the hydroxyl groups on glycerol enable its complexion with transition metal ions (here including Co2+ and Nb5+) to form organic–inorganic complex microsphere particles, which ensure the uniform distribution of Co. With the subsequent detachment of organic component, the porous structure of Co–NbO is naturally generated. To further fabricate PDA–Co–NbO, we performed the in situ polymerization of dopamine hydrochloride on the Co–NbO surface. Then the PDA conformal coating is carbonized to construct the conformal carbon layer on PDA–Co–NbO with better defined porous structure due to the etching effect of PDA. In Figure 2O5 . The pristine T–Nb2O5 (denoted as P–NbO) is also correspondingly synthesized without the addition of Co precursor as a reference. The most intense peaks at 22.6°, 28.3°, 28.9°, and 36.5° are indexed to (001), (180), (200), and (181) planes, respectively, manifesting the sufficient crystallinity for all three samples. Note that no extra peaks of undesirable impurities  are found for Co–NbO and PDA–Co–NbO, indicating the successful lattice doping of Co into T–Nb2O5. The bulk doping effect is also disclosed by the positive shifting of XRD peaks in view of the smaller ion radius of Co2+ (0.72 Å) or Co3+ (0.53 Å) than that (0.78 Å) of Nb5+  and pentagonal bipyramids (NbO7) through shared corners or edges, which is also called as self‐supporting “room and pillar” structure as shown in Figure  It is deduced that the doped cobalt atoms likely occupy the niobium sites in NbO6 and NbO7 polyhedra, without the serious change of overall structure framework , the Nb–O polyhedra in layer are highly distorted, and it complicates the atomic arrangement and cell structure. The partial substitution of cobalt on niobium sites and reduced valence of Co2+ (3d7) mitigate the cation–cation repulsion and influence the electronic states of Nb5+ (4d0). Therefore the existence of lower‐valence Co dopant is expected to cushion the Jahn–Teller effect on the distortion of Nb–O polyhedra, and to stabilize Li‐ion migration channels and contribute to the lattice diffusion kinetics.To further improve the capacity performance, less than 10 mol % Co or Mn is doped into T–Nbon of Co.22 Wit+ Figure . The dop5+, centered at 207.2 eV (Nb5+ 3d5/2) and 209.9 eV (Nb5+ 3d3/2), with a spin–orbit splitting energy of 2.7 eV and an approximate peak area ratio of 3:2. The O 1s XPS spectrum in Figure  The presence of the smaller shoulder peak is attributed to the residual oxygen during the carbonization of polydopamine or the potential bonding between C coating and lattice O in T–Nb2O5. The XPS spectrum of N 1s in Figure  The doping of trace amount of nitrogen and oxygen in PDA derivative carbon layer is beneficial to improve the lithiophilicity of PDA–Co–NbO surface and promote the uptake and adsorption of electrolyte. It is possible that the high adsorption energy of Li on N atoms enables the facile detachment of solvation shell of Li+. No obvious peak is observed in the XPS spectrum of Co 3d in Figure X‐ray photoelectron spectroscopy (XPS) of PDA–Co–NbO was employed to analyze the surface chemical composition. The XPS spectrum of Nb 3d in Figure o of 3:2.26 Theg values of 2.012 and 2.052. Co doping enables the formation of electron‐hole pairs in T–Nb2O5. The formation of oxygen vacancies resulting from cobalt doping can be described in the following formulao represents the lattice oxygen, VO·· represents the oxygen vacancy with two units of positive charge, Co′′′Nb and Co′′Nb denote the substitution of Co atoms in Nb lattice sites with three and two units of negative charge respectively. The superoxide radicals (O2•−) can be formed if the free electrons in electron–hole pairs are captured by oxygen atoms as described in the following formulaElectron paramagnetic resonance (EPR) was employed to prove the existence of cobalt, and the magnetic signal of EPR originates from the unpaired electrons. In Figure nd 2.052.31 Co g = 2.012 should be associated with the existence of superoxide radicals (O2•−). And the peak at g = 2.052 should stem from the unpaired electrons on cobalt, e.g. three unpaired electrons in Co2+ and four unpaired electrons in Co3+. In this case the magnetic signal coming from cobalt likely influences the detection of superoxide radicals, leading to a subtle positive shift of g value. In the EPR curve of PDA–Co–NbO, there are two obvious peaks centered at g = 2.001 and g = 2.013. The elimination of peaks corresponding to cobalt possibly originates from the electromagnetic‐interference shielding of surface carbon layer in PDA–Co–NbO. Both the signals at g = 2.001 and 2.013 should be associated with the stabilization of superoxide radicals but with different coordination environments, but they are much weaker than that at g = 2.012 for Co–NbO. The second annealing for PDA pyrolysis in inert atmosphere is responsible for the healing of excess O2•− and therefore the attenuation of EPR signals. Note that the peak at g = 2.001 is probably assigned to the stabilization of oxygen radicals anchored at the interface with O/N doped carbon coating. Raman spectra were used to explore the vibration situations of metal–oxygen bonds in Co–NbO as shown in Figure –1 correspond to typical bending modes of Nb–O–Nb linkages. After Co doping, the intensities of the two peaks at 120 and 230 cm–1 are obviously weakened, in view of the potential formation of Co–O–Nb linkages and the resultant disordering of bending in T–Nb2O5. The strong and broad peak at 690 cm–1 corresponds to the symmetric stretching mode of NbO6 and NbO7. Note that no additional peaks assigned to the stretching of Co–O bonds is observed in the range of 300–800 cm–1, further verifying the suppression of CoOx phase precipitation and firm substitution of cobalt atoms on niobium lattice sites in Co–NbO, which benefits from the excellent dispersion effect of Co element in organic–inorganic complex precursor. This EPR result indicates that the Co doping enables the stimulation of free electrons, enabling the improvement of bulk electron conductivity in PDA–Co–NbO and Co–NbO.The peak at s (O2•−).32 AndFigure  which effectively retards the fusion of adjacent particles and optimize the porosity of electrode network. The elemental mapping by energy dispersive X‐ray spectroscopy  imaging was used to further explore the detailed structural and morphology information of PDA–Co–NbO in Figure 2O5, respectively. The uniform carbon layer with a thickness of ≈2 nm can be discerned to conformally decorate on the surface of T–Nb2O5 nanoparticles from the HRTEM images as well  images of Co–NbO Figure 3a,b dispFigure 38 whi2O5 . The electrochemical window of T–Nb2O5 was usually confined to 1–3 V in most literatures, which corresponds to two Li+ insertion in one Nb2O5 unit. Here the electrochemical window is extended to 0.4–3.0 V to improve the specific capacity. The corresponding galvanostatic charge/discharge tests of Li/P–NbO, Li/Co–NbO, Li/PDA–Co–NbO and Li/Mn–NbO cells were performed to testify the extra contributions of Co/Mn dopant and conductive carbon layer on cycling performance. Figure –1. PDA–Co–NbO displays the higher specific capacities at all the current densities  than that of Co–NbO (246 mAh g−1). During the long‐term cycling for at least 600 cycles, the coulombic efficiency (CE) plots are highly stabilized at 100%. It is evident that the extra wiring of conformal carbon layer endows PDA–Co–NbO with better cycling durability, which is also supported by the superior long‐term cycling performance at a higher current density of 1 A g−1 in Figure −1 for three cycles). It needs to be emphasized that the Li/PDA–Co–NbO cell can preserve a specific capacity as high as 143 mAh g−1 after cycling for extra‐long 8500 cycles. Within the 1st–2000th cycles, Li/PDA–Co–NbO cell exhibits the obviously better capacity retention (69%) than Li/Co–NbO cells (57%). The corresponding charge/discharge curves of Li/PDA–Co–NbO cell display the high symmetry and excellent reversibility even after cycling for 5000 cycles  with areal capacities of 0.307, 0.419 and 0.668 mAh cm−2 at the areal loadings of 2.0, 3.0, and 4.5 mg cm−2, respectively at the 150th cycle and 500 mA g−1  within 250 cycles. The wiring and infiltration effects of firm conformal carbon layer at the surface of each primary particle of PDA–Co–NbO are responsible for the preservation of high areal capacity even under the high loading and long cycling conditions. The areal capacity advantage of PDA–Co–NbO at different current densities is displayed when comparing with other representative intercalation anodes  in Figure  This areal capacity value of 0.807 mAh cm−2 for PDA–Co–NbO are more than two times of those for other typical anode materials at the same current density of 500 mA g−1.The electrochemical activation process in P–NbO mentioned in Figure 2O5 unit.39 Her1 Figure , the PDA1 Figure . When thFigure −1, corresponding to the charge/discharge platforms in Figure −1, while Li/Co–NbO cells embrace highest potential gap at all rates except 0.1 mV s−1. Figure −1), and the steeper slope reflects the higher diffusion coefficient according to Formula S1 (Supporting Information). Obviously, the peak current responses and corresponding slopes for PDA–Co–NbO and Co–NbO highly exceed those for P–NbO, benefiting from the superior electrochemical kinetics for the former two. The higher diffusion coefficient  is attributed to the synergistic contributions of conductive carbon layer network, NbO domains and cobalt dopants. The lithium diffusion coefficient can be quantitatively estimated by the galvanostatic intermittent titration technique (GITT), based on the following formulaD is the Li diffusion coefficient, S and L refer to the thickness and surface area of electrode material, respectively, mB, Vm, and MB are denoted as the mass loading, theoretical molar volume (50.7 cm3 mol−1) and theoretical molecular weight (265.81 g mol−1) of T–Nb2O5, respectively. As visualized in the GITT curves . The Eτ value is equal to the transient potential during the pulsing step and it displays a linear relationship (R2 = 0.998) with the square root of pulsing time  are roughly two order of magnitude higher than those of fast‐charging type Li4Ti5O12 (6.77 × 10–14 to 2.8 × 10–15 cm2 s−1) reported in literatures . When SOC and SOD are lower than 80%, the D values for PDA–Co–NbO are higher than for Co–NbO and they display an overall downward tendency with the increase of SOC and SOD. These phenomena indicate that both the deep lithiation and delithiation would somewhat degrade the electrochemical kinetics. The Li‐stuffed and Li‐emptied channels likely become less stable or more distorted as observed from the potential conversion into NbO phase and amorphization of lithiated LixNbyOz phase, which would be further discussed below.The calculated Li diffusion coefficients . When the b value is close to 0.5, the current response is controlled by semi‐infinite linear diffusion, and the surface capacitance dominates the current response when the b value approaches 1. As shown in Figure k1v) and pseudocapacitive current (k2v1/2) at a given voltage (V) according to the following equationThe intercalation pseudocapacitance should be also responsible for the high‐rate performance of PDA–Co–NbO and Co–NbO, and it was quantitatively analyzed based on the collected CV curves at different scan rates. The dominant electrochemical step can be primarily judged through the assumed power‐law relationship of peak current versus sweep rate i = avb.[48]k1 and k2 can be obtained through the linear relationship of i(V)/v1/2 versus v1/2. The calculated pseudocapacitive current plots  within the total CV curve can be integrated into the green area as shown in Figure −1), and the corresponding areas for P–NbO and Co–NbO are outlined in Figure −1). This enhancement is in accordance with the better rate performance for the formers, and is ascribed to the coral‐like porous electrode structure (with higher specific area surface), defect and coating reinforced charge transport pathways.Therefore, the parameters −1 in advance to stabilize the solid electrolyte interphase (SEI). The obtained Nyquist plots with frequencies from 1 MHz to 0.01 Hz and at stepwise increasing temperatures (30 – 60 ℃) are depicted in Figure  The typical well‐defined semicircle at high frequencies is assigned to the resistance of solid electrolyte interface (denoted as RSEI), the incomplete semicircle at medium‐frequency range is attributable to the change transfer resistance (denoted as RCT), and the steep tail (nearly straight line) at low frequency zone is simulated with a finite‐length Warburg impedance component (Zflw), which presents Li+ diffusion and accumulation at the electrolyte/electrode interface. The specific frequency ranges of above‐mentioned three regions can be found in Figure RE), and there is almost no difference of RE values  among all the three cells. The Li/PDA–Co–NbO cell enables the smallest resistance contributions from charge transfer and SEI, and therefore the smallest total interface resistance RTotal  was employed to further explore the evolution of electrochemical kinetics after the introduction of cobalt and carbon layer. Before the EIS measurement, the Li/PDA–Co–NbO, Li/Co–NbO and Li/P–NbO cells were cycled for 5 cycles at 100 mA gn Figure 53 The2O5 particles with better electrochemical activity. To clarify the structural evolution, a series of characterizations on the cycled electrodes were adopted as shown in Figure Pm3m). What is more, a broad peak can be observed ranging from 20 to 35°, and it can be ascribed to the amorphization of LixNbyOz , indicating the electrochemical irreversibility of NbO conversion product, which is expected to serve as electronically conductive wires to promote the cycling stability of LixNbyOz. Note the fraction of this nonactive phase is not high and therefore the capacity performance would not be compromised. This result also implies that the primary Nb2O5 would not undergo the deep conversion reaction and the appearance of NbO is likely caused by the shallow phase segregation of Li‐stuffed LixNb2O5 or LixNbyOz into NbO and Li2O. The local phase precipitation is sufficient to alleviate the lattice stress of LixNbyOz and retard its further decomposition. The electrochemical capacity release should stem from the formation of dominant Nb4+ peaks (at 206.8 eV for 3d5/2 and 209.5 eV for 3d3/2) at the discharge state. The Nb4+ species can be reversibly converted into the Nb5+ species (with dominant peaks at 207.3 eV for 3d5/2 and 210.1 eV for 3d3/2) after recharging. Therefore, the lithium storage mainly depends on the reversible conversion between Nb5+ and Nb4+, and the NbO domains mainly assist the migration of electron flow. The partial density of states (PDOS) and corresponding band structures of NbO and T–Nb2O5 were calculated through density function theory (DFT) to unveil the electrical conductivities of NbO and T–Nb2O5 in Figure 2O5 is 1.75 eV. The conductive NbO has the overlapped conduction band (CB) and valence band (VB), since both the CB and VB are composed of Nb 3d orbitals. The calculated PDOS of T–Nb2O5 shows that VB is mainly composed of O 2p and Nb 4p orbitals, while CB is mainly composed of Nb 3d orbitals. This result strongly proves that the cycled electrode of NbO/LixNbyOz displays the obviously higher electrical conductivity than T–Nb2O5. The calculated XRD patterns of NbO and T–Nb2O5 structures used for PDOS calculation are displayed in Figure 2O5 phases. The in situ Raman spectra , do not undergo the position shifting during the whole cycling. The shallow conversion reaction into NbO and irreversible Li consumption by carbon additive below 1 V do not degrade the electrolyte stability.Transition metal dopants  and nitrogen doped carbon layer can synergistically endow the T–Nbre Pm3m).55 Wha Figure .[56]a Figure  of cycleFigure 5+ (0.78 Å) to Nb4+ (0.83 Å). On the other hand, the lattice expansion may also be caused by the lithium occupation on the interstitial sites of NbO7/NbO6 polyhedra. In the following charge process, the highly reversible positive shifting is observed with lithium extraction. The in‐situ XRD evolution indicates that the dominant electrochemical behavior is the reversible solid‐solution‐like (de)lithiation without the obvious appearance of phase transformation during the first cycle. The excellent reversibility of XRD pattern evolution is also illustrated by the corresponding 2D contour plots in Figure +, agreeing with the solid solution behavior  are likely responsible for the acceleration of structure disorder, although it is still based on the solid solution reaction behavior. The stoichiometric loss should be triggered by the phase segregation from Li‐stuffed Li2Nb2O5 with the formation of NbO and Li2O domains. The (001) peak signal can be preserved with less broadening and it does not vanish during the following charge, but it cannot return to the original position before the beginning of the second cycle. Instead, this (001) peak experiences a marginal shifting during the second charge. It appears that the “room and pillar” constructed layers along c axis are more resistant to the structure collapse or disordering. The corresponding lattice parameter undergoes an increase from 3.92 to 4.10 Å with a larger magnitude during the second discharge compared with the case in the first discharge  appear during the second charging process, respectively corresponding to the (110), (111), and (220) planes of NbO in Figure th cycle, as the once‐formed NbO domains would not be electrochemically oxidized and they only serve as electronic wires. It is speculated that the intercalation of a large amount of lithium incurs the irreversible lattice expansion of Nb2O5, leading to the partial structural disorder of T–Nb2O5 with the formation of more crystallized NbO and less crystallized LixNbyOz at the same time, similar to the report of disordered rock salt Li3V2O5 phase as the lithiation product of V2O5.In contrast, in the second discharge process  and Co–NbO/LiFePO4 full cells. After the cycling of Co–NbO at 100 mA g−1 for 5 cycles in advance (with lithium as counter electrode), the SEI‐stabilized Co–NbO was disassembled and adopted as anode in full cells. The typical charge/discharge curves of Co–NbO anode and NCM811 cathode in Figure −1 after 200 cycles with an average decaying rate as small as 0.09% per cycle at a rate of 0.5 C (1C = 280 mA g−1) at room temperature. The Co–NbO/NCM811 full cell delivers the reversible capacities of 116.3, 103.6, 88.7, 63.0 mAh g−1 at the rates of 0.2, 0.5, 1.0, and 2.0 C, respectively under a smaller N/P ratio of 0.5  and high‐rate performance (144.1 mAh g−1 at 10.0 A g−1). The multiple wiring effect also enables a high mass loading of 4.5 mg cm−2 and high areal capacity of 0.668 mAh cm−2 (even after 150 cycles) for PDA–Co–NbO. The exploration of more conductive wiring models in T–Nb2O5 is a potential solution to its practical application in high‐rate LIBs.In summary, we propose a triple conductive wiring strategy through electron doping, chelation coating and electrochemical conversion inside the micro‐sized aggregates of T–NbThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


Our approach makes use of the entropic advantage of studying a weak force intramolecularly, embedded into a cooperatively folding system, and of the combined use of NOE‐ and RDC‐based ensemble analyses to accurately describe the orientation of the donor and acceptor sites. Thermodynamic constants , describing the specific interaction, were derived from variable temperature chemical shift analysis. We present a methodology for the experimental investigation of remarkably weak halogen bonds and other related weak forces in solution, paving the way for their improved understanding and strategic use in chemistry and biology.Detection and characterisation of very weak, non‐covalent interactions in solution is inherently challenging. Low affinity, short complex lifetime and a constant battle against entropy brings even the most sensitive spectroscopic methods to their knees. Herein we introduce a strategy for the accurate experimental description of weak chemical forces in solution. Its scope is demonstrated by the detailed geometric and thermodynamic characterisation of the weak halogen bond of a non‐fluorinated aryl iodide and an ether oxygen (0.6 kJ mol Very weak halogen bonds are excessively difficult to characterise in solution. We present an approach for the detailed NMR spectroscopic description of very weak interactions, providing the solution ensemble of the flexible model compound, the accurate geometry of the binding sites and the corresponding thermodynamic parameters. Here NOE abbreviates nuclear Overhauser effect, RDC corresponds to residual dipolar coupling, and ΔT indicates variable temperature experiments. A halogen bond (XB) is the attractive interaction of a polarized halogen and a Lewis base.Intramolecular assessment of weak interactions offers an entropic advantage, and accordingly, the halogen bond of a perfluorinated donor could be assessed even in polar solvents, upon incorporation into an intramolecular model system.Herein, we present a strategy for the remarkably accurate spectroscopic characterisation of weak interactions in solution. Making use of a non‐cyclic, cooperatively folding backbone and the entropic advantage of an intramolecular setting, the method allows the thorough analysis of one specific weak interaction. We introduce the use of orienting media to the research field of halogen bonding, for the detection of bond vectors upon observation of residual dipolar couplings (RDCs). This allows the experimental detection of the orientation of halogen bond donor and acceptor sites, and thereby offers a direct evidence for the formation of a halogen bond in solution, without enforcement of the bonding geometry. The combined use of NOE‐based ensemble analysis and RDCs of a highly flexible molecular system is novel. The scope of this strategy is demonstrated by the detailed characterisation of an iodine‐centered, weak halogen bond in solution that has previously been inaccessible, being too weak to be detected by current techniques. Providing a new strategy to study the weakest interactions in solutions, we expect to facilitate future developments and applications of halogen bonding in, for example, medicinal chemistry and catalysis.1, Figure 2). In the latter, the formation of a halogen bond is prevented by the smallest possible change, an −O− to −CH2− substitution, to eliminate the halogen bond acceptor site. This reference molecule, 2, resembles the halogen bonding model system in all respect, but the ability to form a halogen bond. The difference in the folding‐defolding equilibrium of the two compounds, 1 and 2, originates from this structural difference. All other interactions contributing to the cooperative folding are identical. Hence, the difference in their folding reflects the thermodynamics of the halogen bond of interest. DFT model computations corroborate that incorporation of the I…O halogen bond into model system 1 does not significantly alter the interaction energy . As the side chains of Thr2, Glu4, Lys7 and Ser9 orient to the opposite face of 1 as compared to those of Phe(I)3 and Hse(Me)8, these do not influence the studied halogen bond.The solution NMR spectroscopic characterisation of halogen bonds1 and 2 were synthesized on solid‐phase and purified by HPLC; for details see the Supporting Information.Compounds 1 and 2, each in 10 K steps over a 130 K temperature range. To allow detection in such a wide temperature range (218 K to 348 K) and in a comparably polar environment, a solvent mixture of CD2Cl2:DMSO‐d6 (4 : 1) was used. The chemical shifts of the completely folded and completely unfolded state were estimated, following the procedure described by Munekata,Thermodynamic analysis was performed by observing the temperature dependence of the chemical shifts of ten backbone protons for compounds obs is the observed chemical shift at temperature T, δU and δF are the chemical shifts at completely unfolded and folded state. Correspondingly, ▵Hm denotes the enthalpy change upon unfolding at the melting temperature, Tm, whereas R is the molar gas constant. To determine the overall folding rate, we normalized the individual protons’ folding curves to the maximum shift change of each. Subsequently, the melting curves, describing the overall folding of the two systems ), decreasing the uncertainty of the estimation as compared to determination of the absolute thermodynamic constants:providing a ΔΔobs is the measured chemical shift, δU‾and δF‾are the shifts in the completely unfolded and completely folded states of compounds 1 and 2, respectively. KF1/2can be deduced as the slope when plotting δobs1-δF1‾δU2‾-δobs2against δU1‾-δobs1δobs2-δF2‾ is the melting temperature of 2 and KF1/2is the ratio of folding constants, which is in agreement with ΔΔG° estimated using the method of Munekata  were determined from spectra acquired on a 3 mM solution of 1 and 2, respectively, in CD2Cl2/DMSO‐d6 (4 : 1) on a 500 MHz NMR spectrometer equipped with a cryogenic probe. For highest accuracy, NOESY build‐ups with seven mixing times, 100–700 ms, were acquired, and inter‐proton distances derived using the initial rate approximation.3JCαH,NH coupling constants were deduced from 1H NMR spectra. The population‐averaged experimental data  were deconvoluted into population‐weighted ensembles of solution conformations using NAMFIS.−1 energy window from the global minimum. Combination of conformers generated by several Monte Carlo searches applying different force fields ensured sampling of the entire conformational space. As the force fields implemented into the software Schrödinger are not parametrized for halogen bonding, we enriched the theoretical input ensemble of 1 with conformers, in which the I⋅⋅⋅O halogen bond was facilitated by an additional force constant promoting an I⋅⋅⋅O interatomic distance shorter than the sum of the van der Waals radii of I and O. Following this procedure, conformational pools of 236 and 176 conformers for 1 and 2, respectively, have been used as theoretical inputs for the NAMFIS algorithm, which thereof identified 9 conformers for 1 and 11 for 2. As the wide conformational variability of the side chains is not possible to cover accurately with the current computational techniques, reliable conclusions can only be drawn regarding the conformation of the backbone. The individual conformations were assessed on the formation of a type II’ β‐turn, on the average interstrand backbone distance and on dihedral angles based on Ramachandran plot analysis to identify folded and unfolded geometries . We confirmed the alignment by detection of a clean quadrupolar splitting (ΔνQ∼125 Hz) of the DMSO‐d6 deuterium signal. Residual dipolar couplings (1CHD) were deduced as the difference of the total couplings (1TCH) measured in the anisotropic and the scalar coupling (1JCH) obtained in the isotropic solution. They were ‐68 Hz to 58 Hz in magnitude  and condition number SVD=4.028. Here, the quality factor (Q) is the measure of the goodness‐of‐fit of the theoretical model (structure) to the experimental data, with Q<30 % indicating a good fit.2‐penalty function, as implemented in the Fitter tool of the software MSpin. Residual dipolar coupling analysis indicated 53 % population of folded conformers for 1, which is in excellent agreement with the outcome of the NAMFIS ensemble analysis (56 %) and with the thermodynamic data obtained from the assessment of the chemical shift melting curves. It further revealed that 1 adopts a folded β‐hairpin encompassing an interstrand I⋅⋅⋅O halogen bond with 30 % population. This corroborates our interpretation of the NAMFIS and variable‐temperature chemical shift analyses, i. e. that the higher folded population of 1 in respect to 2 is due to halogen bonding. Singular value decomposition selected β‐hairpin conformations with a C−I⋅⋅⋅O angle of 145°, from input conformations possessing a wide variety of donor‐acceptor orientations. This observation is in excellent agreement with the bimodal angle preference of ∼160–170° and ∼145–150° for halogen bonds in complex protein‐like environments, as described by Ho et al.Whereas NAMFIS analysis provides information on the overall fold, it does not allow conclusions to be drawn on the orientation of the functionalities involved in halogen bond formation. To assess the latter, we collected residual dipolar couplings for uplings 1DCH were−1). This was achieved by incorporation of the halogen bond donor and acceptor sites into a cooperatively folding system that allowed characterisation of its bond strength and orientation as an intramolecular interaction. In order to demonstrate the impact of the halogen bond on cooperative folding, a reference system lacking the halogen bond acceptor site showed, by NOE and scalar coupling based NAMFIS analysis, to fold 17 % less, as compared to the system designed to form an intramolecular halogen bond. Out of the 56 % folded conformers of the latter, 30 % formed a halogen bond, as shown by the analysis of 1H,13C residual dipolar couplings obtained by using a dilute solution of alignment medium (PBLG). Importantly, the detection of residual dipolar couplings allowed identification of the C−I⋅⋅⋅O halogen bond angle, 145°, which has previously not been conceivable in solution. It should be noted that this I⋅⋅⋅O halogen bond is much too weak to be detected in solution in an intermolecular setting using standard NMR techniques, whereas the strategy presented here allowed its detailed energetic and geometric characterisation.We disclose a strategy for the thorough geometric and thermodynamic characterisation of very weak interactions in solution, demonstrating its scope by the description of an I⋅⋅⋅O halogen bond  should be addressed to the authors.Supporting InformationClick here for additional data file.


We show that the formation of a perovskite structure type compound with the general formula ABO3 (where A and B are cations and O is an oxide ion) becomes energetically favorable when it has a higher Madelung electrostatic potential than the combined Madelung electrostatic potential of parent binary compounds AO and B2O3 or BO2. It is further shown that strong lattice site potential results in stability of high valence or high valence ions can be stabilized in a lattice site with strong lattice-site potential. It further follows that certain ions experience maximum lattice site potential at the B ion lattice site of the perovskite structure when compared to other structures such as fluorite BO2, rutile BO2 and corundum B2O3. Therefore, (i) the stability of an ion with a high (and uncommon) valence state at the B site being higher than that at the A site, (ii) occurrence of point defects at A or O sites with weak lattice site potentials, respectively and (iii) instability of perovskite A4+B2+O3, and A5+B1+O3 compounds, respectively can be rationalized by lattice site potential and Madelung electrostatic potential analysis.Valence stability of aliovalent ions is mostly correlated with lattice site potential in ionic crystals. Madelung electrostatic potential is obtained by adding all the lattice site potentials for all the ions present in a crystal structure. Therefore, valence stability and the stability of a crystal structure can be better understood with consideration of both the lattice site potential and Madelung electrostatic potential. This was first demonstrated more than four decades ago by one of the present authors. We revisit this situation by using re-calculated lattice site potential and Madelung electrostatic potential for perovskite structure type ABO The stability of perovskite structure type oxides and higth valence state of cations, respectively may be better understood by considering lattice site potential and Madelung electrostatic potential energy. Metal oxides with chemical formula ABO3 are a subclass of this large family of compounds. For many years, perovskite oxides have been studied extensively in the literature of chemistry, physics and geology for a large variety of solid state phenomena associated with electric, optical, thermal and magnetic properties. A large number of oxides with perovskite type structures are possible to synthesise in a laboratory for applications in vast areas spanning from electronics to catalysis. It is of immense importance to understand the factors that may correlate with the formation and stability of perovskite structure type ABO3 compounds for the design of new compounds with novel structures containing desired valence states of B cations.An enormous number of solids show crystallographic structural similarity to the mineral perovskite  and was considered for the first time in 1974.8 This was referred by Rustam Roy in 1975 in a plenary lecture presented as the Robert B. Sosman Memorial Lecture at the 77th annual meeting of the American ceramics society and was published in 1977.9 This is a novel and very less explored concept. We also note that not many researchers have been able to consider this concept perhaps due to the inaccessibility of the original article.8 Here, the present article is written by re-visiting Madelung lattice energy and site potentials based upon recent data calculated using VESTA10 program and recent crystallographic  and thermodynamic data. We emphasize on the fact that ionic solids are stabilized due to the gain in lattice energy that is reflected in Madelung electrostatic potential (EM). Nevertheless, recently, Hoppe demonstrated that the enthalpy of formation of solids is reflected in Madelung electrostatic potential in 199511 and Glasser has shown that for ionic solids the lattice energy and Madelung electrostatic potential energy (EM) has a linear relationship in 2012.12The reason for why many compounds would take perovskite structure type ABX3 compound has a cubic unit cell. This structure can be constructed from corner shared BO6 octahedra with A cation being present at the 12 coordinated interstitial position within the octahedral network  the presence of large A cation within the interstitial position, (ii) crystal field splitting and (iii) hybridization of orbitals.2The ideal perovskite structure type of ABO network . A and BEM is Madelung electrostatic potential energy and EN is combination of all other energies .13The total internal energy of ionic solids may be expressed as3 compounds (with varying combination of A and B) led to a set of fundamental criteria for the formation of perovskite structure type ABO3 with varying combination and valence states of A and B and summarized by Yoshimura et al.8,14 as follows: oxides of A can be reacted with oxides of B to form perovskite structure type ABO3 compounds if: (i) B ion is smaller than A ion, (ii) valence of B ion is larger or equal to A ion and (iii) sum of valence of A and B ion is six . An in depth analysis showed that the valence stability originates from gain of lattice site potential. Recently, there have been some excellent research work reported in the literature based on the use of Madelung electrostatic potentials energy for understanding the stability of perovskite structure type compounds. Shan et al. investigated Li ion insertion in perovskite structure type SrVOδ3−, La2/3TiOδ3− and TiOδ3−, from the lattice site potential analysis and found that the insertion of lithium ion is limited by the effect of lattice site potential on the valence stability of V and Ti.15 In these perovskite structure type oxides, the valence state of tetravalent V and Ti can be reduced to 3 upon insertion of monovalent Li ion at an interstitial site. It is interesting to note that for spinel structure type LiMn2O4 compounds, Ragavendran et al. have shown that the Madelung electrostatic potential of LixMn2O4 almost linearly correlates with x and therefore voltage of Li-ion battery made using LiMn2O4 can be easily estimated from the Madelung electrostatic potential.16 Furthermore, the perovskite structure-type is not limited to ABO3 and a large number of organic–inorganic halide crystallise into perovskite type structure with general formula ABX3. A bottle neck for the realization of full potential of so called ‘perovskite solar cells’ and related electronic devices has been the degradation of electronic devices fabricated using hybrid organic–inorganic ABX3 compounds.17 The degradation of these devices arise from instability of the perovskite structure type hybrid compounds and can also be understood from Madelung electrostatic potential analysis.18 Furthermore, it has also been shown that the stability of oxides such as Li4SiO4 towards the reaction Li4SiO4 + CO2 → Li2SiO3 + Li2CO3 has direct consequence on the CO2 capture and storage abilities and recent study by Oh-ishi et al. had successfully employed lattice site potentials analysis of all Li+ ions in the unit cells of Li4SiO4, Li2SiO3, and Li2CO3 to evaluate the extraordinary stability.19The analysis of Madelung electrostatic potential and lattice site potential of a large number ABO3 oxides using an ionic model developed by Zaanen et al.20 and demonstrated in a large number of oxides by Torrance et al.21 and Arima et al.22 Kato et al. employed lattice site potential analysis to show that the valence band structure can be tailored by designing the stacking sequence of layers of a number of layered compounds BiOX , Bi4NbO8X , Bi2GdO4X , and SrBiO2X  and therefore, lattice site potential analysis may be useful for designing new photocatalysts.23 Here we look back at the development of Madelung electrostatic potential and lattice site potential only for perovskite type ABO3 oxides. We show that the lattice and valence stabilities of perovskite structure type ABO3 compounds can be understood from simple Madelung electrostatic potential view point and valence stability is directly related to the lattice site potential. Based on the existing literatures, our hypothesis is that the stability and properties of perovskite type ABO3 oxides can be rationalized better in terms of change in Madelung electrostatic potential energies rather than structural perturbation such as distortions and tolerance factors.24,25On the other hand, lattice site potential can be used for the estimation of the oxygen to metal charge transfer and bandgap of perovskite type ABO10 to recalculate the lattice site potentials and Madelung electrostatic potential directly from the known crystal structures. The Madelung electrostatic potential EM obtained from VESTA has the unit of MJ mol−1 asymmetric unit. In order to compare EM per mole of each compound with different crystal symmetry (space groups), the EM output of each compound from VESTA was multiplied by the multiplicity of general Wyckoff position and divided by the number of formula units in the unit cell.The Madelung electrostatic potential and lattice site potentials can be calculated using commercial first principle calculation software packages. We have used freely available software VESTA6,7 the Madelung electrostatic potential energy EM can be expressed asqj is total charge number of lattice point j, pj is frequency of occurrence of j lattice point in the unit cell, ϕj is the lattice site potential (Å−1) at the lattice point j, and κ is total number of molecules in the unit cell.According to Van Gool and Piken,Ma is expressed in terms of the characteristic length a, such as the lattice parameter of unit cell. For compounds with complex structures an arbitrary characteristic length r such as nearest neighbor atomic distance can be defined and the Madelung constant can be expressed asThe Madelung constant can be expressed asFrom 10 lattice site potential ϕi and Madelung electrostatic potential EM of a given structure is calculated using following equationsZj is valence of ion j, ε0 is permittivity of vacuum and lij is distance between ions i and j.In VESTA,e.g. NaCl, CsCl), the lattice site potential ϕj for each site in the crystal structure is inversely proportional to the characteristic length. However, if the characteristic length is fixed, for the first time Yoshimura et al.8 had analyzed that, lattice site potential ϕj, the Madelung constant Ma and Madelung electrostatic potential energy EM varies with the valence states. ϕj as well as Madelung electrostatic potential EM in ideal cubic perovskite type structure of ABO3 with fixed lattice constant of 3.881 Å. Based on Ma also shows similar trend with valence. ϕA) and O site (ϕO) increase with the valence (0–6), (1–5), (2–4), (3–3)  while the lattice site potential of B ion (ϕB) decreases. This re-calculated trend of lattice potential with valence pairs is in excellent agreement with that obtained from method developed by Van Gool and Piken,6,7 and adapted by Yoshimura et al. in 1974.8 All calculations were performed using VESTA10 employing the radius of ionic sphere and reciprocal space range of 1.6 Å and 10 Å−1, respectively. For the calculation of ϕ and EM of A0B6+O3, a dummy A such as Na atom was placed at (½ ½ ½) and B cation Re6+ at (0 0 0) atomic coordinates, respectively of space group Pm3̄m (221) with both valence and occupancy, respectively of A were set to zero. The calculation for valence pairs (4–2) and (5–1) assumed hypothetical A4+B2+O3 and A5+B1+O3 perovskite structure type compounds because they rarely exist.For similar structures with fixed atomic position (ϕ) of each ionic site and Madelung electrostatic potential (EM) energy changes with lattice parameter when valence is fixed, for each valence pair in each perovskite structure type ABO3. The variation observed in et al. in 2000  of the decomposition reaction BO2 → ½B2O3 + ¼O2 and corresponding Born–Haber cycle , one can postulate that the stability of high valence ions is enhanced in crystal structure with large lattice energies. A similar analogy can be drawn for the valence stability in complex oxides where multiple ions are present such as in perovskite type ABO3 compounds. Furthermore, cations provide large electrostatic potential contribution to the total Madelung electrostatic potential in crystal structure as shown in the plots of ϕA and ϕB in et al.24 Therefore, the valence stability of cation is rather directly related to the lattice-site potential of the cation site in a crystal structure and the stability enhances when they are present in lattice site with strong lattice site potential such as B site in perovskite structure type ABO3 oxides.2+B4+O3 oxides but not in fluorite type B4+O2. For example, investigation on the synthesis of perovskite type ABO3 compounds had revealed that several stable perovskite type A2+B4+O3 compounds  can be synthesized and are stable at room temperatures in spite of the poor ambient condition stability of Pr4+O2 and Tb4+O2 and existence of both Ce23+O3 and Ce4+O2.27 In order to explain this experimental fact, an in depth analysis of lattice site potential and Madelung site potential was presented in.8 Here, we have recalculated the lattice site potential and Madelung electrostatic potential using VESTA10 and presented in 8 In 3 which we believe due to not so accurate estimation of lattice parameter for TbO2.A number of experimental study has indicated that high valence of some B cation is stable in perovskite type A4+ was not stable in fluorite structure type B4+O2 but stable in perovskite structure type A2+B4+O3 is in accordance with the data presented in 3 in each valence state situation. Therefore, we find that the formation of these perovskite type oxides are associated with significant gain in Madelung electrostatic potential (ΔEM) calculated by subtracting the combined EM of rock salt structure type AO and fluorite structure type BO2 from the EM of perovskite crystal structure type of final ABO3  for the reaction A2+O + B4+O2 → A2+B4+O3 of several oxides .13 They also calculated change in Madelung electrostatic potential (ΔEM) from the EM of A2+B4+O3, AO and BO2 and showed that heat of formation of perovskite structure type compounds can be written as EM correlates well with Goldschmidt tolerance factor rA, rB and rO are ionic radius of perovskite structure type ABO3) and ΔEM becomes more negative when t reduces from 1 to 0.8. Kamata et al. found that lattice site potential for B and Madelung electrostatic potential of A2+B4+O3 , varies linearly with pseudocubic lattice constant and the lattice site potential of Mo is much deeper in perovskite type structure than MoO2.28 It is further interesting to note that some of these perovskite-type oxides are so stable that, they are not easily reduced under hydrogen atmosphere at moderate temperatures. From above examples, it can be reasonably hypothesized that the gain in lattice site potential of B site and hence the Madelung electrostatic potential can promote the formation of perovskite structure type ABO3 compounds. Furthermore, from the lattice site potential consideration, it can be shown that Ca2+Mn4+O3 and A2+Pb4+O3  are stable under conditions where the corresponding MnO2 or PbO2 decomposes to lower valence oxides.8Takayama-Muromachi 2+B4+O3 can be further extended to A3+B3+O3. For example, hexagonal structured A2O3 (such as La2O3) can be reacted with rhombohedral (corundum) structured B2O3 (such as Cr2O3) or rutile structured BO2 (such as CrO2) to form perovskite structure type A3+B3+O3 compounds. However, high valence Ni3+ and Co3+ in corundum structure Ni2O3 and Co2O3 respectively are not stable but LaNi3+O3, LaCo3+O3 are stable. Lattice parameters and atomic coordinates for corundum structure Ni2O3 and Co2O3 is unavailable. In this situation the Madelung electrostatic potential and lattice site potential, can be calculated using a different characteristic length r = rB + rO for each known A2O3, B2O3, BO, BO2, and ABO3 using 8 Further detailed discussion on the characteristic length and Madelung electrostatic potential can be found in.8,12 However, for comparison purposes, here all lattice site potentials were obtained using VESTA10 that uses 3+O3  and B23+O3 , are given in 3+B3+O3. It can be clearly seen that there is significant enhancement of lattice site potential for B site for all La3+B3+O3 from that of corundum structure B23+O3. Thus, high valence state of Ni and Co is stable in La3+B3+O3 is due to this large lattice site potential at B site and consistent with similar high valence stability situation of A2+B4+O3 as discussed in the previous section. Furthermore, from 2O3 is more difficult to form oxygen related point defects than perovskite structure type La3+B3+O3. Indeed, a similar analysis has been performed for rutile BO2 structure as-well8 and therefore it may be understandable that why oxygen vacancy is rarely seen in rutile or corundum structure although plenty of vacancy present in perovskite structure type A3+B3+O3 compounds. Therefore, lattice site potential analysis provides reasonable understanding of the defect formation in perovskite structure ABO3 compounds and preparation of many perovskite structure type ABO3 phases from constituent binary oxides could be understood due to the strong B site potential and gain of Madelung electrostatic potential ΔEM.The valence stabilization due to strong lattice potential at B site as described above for A4+B2+O3 with valence pair (4–2) cannot be prepared and rarely exist. This can be rationalized from Madelung electrostatic potential view point as well. For example, the calculation of ΔEM of a hypothetical chemical reaction ThO2 + MgO → Th4+Mg2+O3 revealed that a net Madelung electrostatic potential loss of 184 kcal mol−1  would occur for the formation of perovskite structure type Th4+Mg2+O3.14 Thus this reaction will not proceed because there is no Madelung electrostatic potential or lattice energy gain in this perovskite type A4+B2+O3 compound and the same hypothesis is applicable to A5+B1+O3. We also note that very recently A4+B2+O3 compounds Bi4+Ni2+O3 and Pb4+Ni2+O3 have been reported to be made only under very high pressure conditions and these can be considered as very special valence combinations where stability is perhaps favored due to other energy factors such as covalency due to hybridization.29,30The perovskite structure type compound AEM in the preparation of perovskite structure type EuTiO3 from EuO or Eu2O3 and TiO2 or Ti2O3. The calculated ΔEM in 2+Ti4+O3 will be more stable than Eu3+Ti3+O3 and usually any preparation will lead to the formation of Eu2+Ti4+O3. This analysis may be expanded to other compounds which are of importance for stabilizing Eu2+ at specific ionic sites for the development of visible light emitting phosphor materials.In 3. With the use of lattice site potential and Madelung electrostatic potential we have been able to demonstrate that (1) perovskite structure type ABO3 compounds can be synthesized from their parent binary oxides because of gain in Madelung electrostatic potential which directly correlates with lattice energy, (2) Madelung electrostatic potential and lattice site potentials dictates the stability of perovskite structure and valence stability of cation in perovskite structure is direct consequence of strong cation lattice site potential and (3) the formation of defects such as oxygen vacancy or excess is driven by intrinsic lower lattice site potential of oxygen site. These results (1)–(3) can be seen various experimental data13,32–34 however, have not been explained in terms of lattice site potential by the authors. Thus, the present article will contribute to understanding of synthesis, structures and properties of functional perovskite type oxides for various areas of applications.In conclusion, we have revisited some fundamental aspects of synthesis, stability and properties of perovskite structure type ABOThere are no conflicts to declare.


In the light of the presented findings, the latter remains experimentally unproven and computationally questionable.Very recently, Boronski et al. reported the first thorium–thorium bond in a crystalline cluster prepared under normal experimental conditions. By using a range of experimental techniques and computational models, the authors found that the isolated actinide cluster contains at its heart two paired electrons delocalized over the tri‐thorium ring. The recorded Raman spectrum allegedly confirmed the existence of a σ‐aromatic three‐center two‐electron bond. In the following we demonstrate that the experimentally observed broad inelastic scattering bands between 60 and 135 cm 2)3 rather than the σ‐aromatic Th3 bonding. The delocalized (ThCl2)3 bond extends the range of the charge‐shift bonding beyond transition metals to a record seventh row of the periodic table.High symmetry and unusual thermodynamic stability of the recently prepared and isolated crystalline tri‐thorium cluster is determined by the unique multicenter charge‐shift bonding pattern (ThCl This type of delocalized σ‐bonding is very often referred to as σ‐aromatic bonding, and therefore this discovery has been acclaimed as extending the range of the aromatic stabilization effect to a record sixth principal atomic quantum shell and to the seventh row of the periodic table.The experimental realization of actinide‐actinide bonding in isolable molecules has been one of the main targets of synthetic actinide chemistry for decades.mmetric D3h tri‐t3, the authors computationally investigated several model clusters, including [{Th(η8‐C8H8)(μ‐Cl)2}3K2] (3′′) with the Highest Occupied Molecular Orbital (HOMO) reminiscent of a 3c–2e bonding, and its dication [{Th(η8‐C8H8)(μ‐Cl)2}3K2]2+ (3*) without the 3c–2e HOMO. The authors found that the geometric structure of 3′′ better fits the X‐ray diffraction data and the calculated vibration modes match well with the intensive signals observed in the experimental Raman spectrum, which they interpreted as the 3 bonding in 3”“experimental confirmation of the Th.3*, which in fact lacks any direct actinide–actinide bonding . In particular, as shown in Figure −1 in 3′′ can easily be found respectively at 63.6, 74.8 and 102.2 cm−1 in 3*, although the lack of delocalized and highly polarizable HOMO in the latter makes them hardly detectable in the Raman spectroscopy. Furthermore, the displacement vectors on thorium and chlorine atoms in both model clusters have comparable magnitudes, which strongly suggests that the observed broad inelastic scattering bands between 60 and 135 cm−1 represent collective stretching and bending modes of the actinide–halogen rather than actinide–actinide bonding.To understand the electronic structure of 3Cl6 cage becomes particularly evident if one compares the overall Th3 bond order in 3′′, that is 3×0.373=1.119, with twelve polarized thorium–chlorine bonds, 12×0.497=5.964; here the full bond order of 0.497 is a sum of the covalent component of 0.198 and the electrovalent (ionic) component of 0.299 estimated from the Evarestov–Veryazov equation.3Cl6 cage in the model cluster 3′′ are associated with the polarized thorium–chlorine bonds. The topological analysis of the one‐electron density calculated for the model cluster 3′′ fully confirms the existence of twelve polarized thorium–chlorine bonds  does not necessarily imply bonding character as it may result from the overlap of the extremely sized 6d orbitals being “squeezed” inside the Th3 unit due to the highly symmetric (Dh3) arrangement of the thorium–chlorine bonds in the Th3Cl6 cage. In other words, the uniform charge distribution over three thorium atoms in 3′′ may have nothing to do with neither 3c–2e bonding nor the aromatic stabilization. Indeed, the additional analyses of the effect of the 3c–2e HOMO on thermodynamic stability  and orbital energy levels of the remaining occupied molecular orbitals  only further confirm that the existence of the tri‐thorium bonding in the crystalline cluster 3 is computationally questionable.The dominant role of the Th−Cl bonding in the Th3 unit in 3′′ is marginal, which is exactly what one would expect for either an ordinary non‐aromatic ring or a non‐bonded system.It should be mentioned that the reported by the original authors remarkable negative values of the nucleus‐independent chemical shifts (NICS),3. To elucidate the nature of bonding in the Th3Cl6 cage, different orbital conjugation topologies has been considered involving the atomic orbitals 6d (thorium) and 3p (chlorine). Two configurations, in which the in‐phase orbital overlap is particularly effective are presented in Figure 3Cl6 cage clearly act as divalent and monovalent elements, respectively, which again perfectly corresponds to the resonance Lewis structures representing the delocalized actinide–halogen charge‐shift bonding (ThCl2)3  method has been used,3Cl6 cage that in 78 % is determined by the multicenter charge‐shift bonding (ThCl2)3. Interestingly, the characteristic shape of the EDDB function in the Th3 core reveals noticeable reduction of the electron density in the centroid and the bond midpoints, which explains the lack of the ring and bond critical points and reaffirms again that the charge delocalization over the thorium atoms may have nothing to do with neither the three‐center σ‐bonding nor the σ‐aromatic stabilization.The experimental Raman spectrum recorded by Lidle and co‐workers and the computational data and arguments presented above indicate that the unusual actinide–halogen bonding pattern is a vital factor determining the high symmetry and stability of the actinide cluster 3 Figure c. To conDh3) and unusual thermodynamic stability of the crystalline tri‐thorium cluster 3 is mostly determined by the unprecedented multicenter charge‐shift bonding (ThCl2)3 rather than the σ‐aromatic Th3 bond. In the light of the presented findings, the existence of the latter remains experimentally unproven and computationally questionable, and if the tri‐thorium bonding exists at all, it should not be expected to be stronger than the extremely weak Th−Th bonds already reported in the literature.2)3 found in the crystalline cluster 3 extends the range of the charge‐shift bonding beyond transition metals to a record seventh row of the periodic table. This discovery may have broader implications for understanding the chemistry of actinides and future attempts to design and synthesize new stable actinide complexes.To summarize, the presented analysis clarifies inaccurate conclusions of the original study by Boronski et al.,The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.Supporting InformationClick here for additional data file.


This intermetallic molecule [ErIII(ReICp2)3] (ErRe3) starts a family of molecular nanomagnets (MNM) with unsupported Ln-TM bonds and paves the way towards molecular intermetallics with strong direct magnetic exchange interactions—a promising route towards high-performance single-molecule magnets.Magnetic molecules known as molecular nanomagnets (MNMs) may be the key to ultra-high density data storage. Thus, novel strategies on how to design MNMs are desirable. Here, inspired by the hexagonal structure of the hardest intermetallic magnet SmCo III center coordinated solely by three transition metals in a trigonal planar fashion.Molecular nanomagnets exhibit magnetic memory effect originating from each individual molecule. Here, the authors report an intermetallic nanomagnet based on a single Er Their commercial success results from the combination of the strong magnetic anisotropy of rare-earth ions and their direct magnetic coupling with transition metal ions. At the other far end of the current search for novel magnetic materials are magnetic molecules, called molecular nanomagnets5 (MNMs) or single-molecule magnets (SMMs). MNMs are molecule-sized objects with magnetic memory effects governed by quantum mechanics10. They are strongly believed to revolutionize magnetic information storage, but at the moment do not have any real-world applications, as their performance is limited by low blocking temperatures and quantum tunneling of magnetization (QTM)11. Current trends in the design of MNMs13 were ignited by the discovery of terbium double-decker [TbPc2]− 14 and were fueled by the seminal perspective of Rinehart and Long15. Design principles focus on the precise control of the coordination sphere of a single lanthanide center to maximize the easy-axis magnetic anisotropy of the complex and to limit the spin-lattice relaxation rates19. While this approach led to several breakthroughs, including the observation of the magnetic hysteresis loop of molecular origin above the liquid nitrogen barrier for dysprosocenium7 and magnetic field control of the magnetization blocking barrier20, it appears to be reaching its limit as the proposed modifications were not as successful as expected22.Rare-earth intermetallic magnets—samarium-cobalt (SmCo)24. An MNM of this type stands behind the success story of the whole field as it was indeed started by the discovery of a dodecanuclear carboxylate-bridged cluster Mn12—a high-spin molecule (S = 10) comprising eight MnIII and four MnIV ions interacting magnetically via the bridging carboxylate ligands27. However, the indirect character of these magnetic interactions, based on a weak superexchange coupling mechanism, limits the performance of Mn12 to very low temperatures, below the boiling point of liquid helium. Attempts to adapt this particular strategy to dysprosocenium complexes suffer from a similar problem –weak superexchange interactions of the lanthanides with other metal centers through diamagnetic bridging ligands limit their influence on the slow magnetic relaxation30. Therefore, the route towards the high-performance MNMs based on superexchange interactions has been largely abandoned.The original approach to MNMs was based on designing multinuclear systems with strong magnetic superexchange, where several paramagnetic centers are glued together by bridging ligands to form a high-spin molecule. If magnetic anisotropy is present in such a system, it shows MNM properties12, its ultimate goal—the room temperature molecular nanomagnet (RT-MNM)—remains elusive. However, the aforementioned examples of the rare-earth intermetallics: SmCo51 and Nd2Fe14B2 provide clues for the possible direction in the design of RT-MNMs. Focusing the efforts on molecules comprising a rare-earth metal center coordinated solely by transition metal ions34, mimicking the first coordination sphere of the Sm center in the SmCo5 magnet 2Dy2I3 where the direct exchange is operational36. Interestingly, even much weaker indirect interactions were reported to be quite efficient in this matter37.Overall, despite the huge progress in the field initiated by Mnnet Fig.  could en34. Later, it was proposed by Rinehart and Long15 as a possible strategy for molecular nanomagnets and put to use by Nippe et al. reporting MNMs with unsupported direct bonds between the dysprosium ion and 4d (Fe) or 4d (Ru) transition metal ions38. However, the magnetic memory effect (magnetic hysteresis) has not been observed, most probably due to the unfavorable ligand field geometry38 or the unfortunate choice of the rare-earth metal32. The introduction of p-block heavy metals directly into the coordination sphere of the lanthanide was also pursued resulting in interesting examples of MNMs40. Other systems involving direct bonds between actinides and d-metal elements were also reported, but slow magnetization dynamics were not studied43. Noteworthy, the concept of placing transition metal ions in the proximity of Ln-based MNMs employing suitable bridging ligands was pursued for tens of years with several examples showing enhancement of the slow relaxation of the magnetization due to the presence of the transition metal46. However, this strategy, while providing very good control of the connectivity of the target multi-metallic compounds, cannot warrant strong coupling between the two metal ions due to the weak overlap of the relevant orbitals of the lanthanide and the bridging moiety. Only the direct bonding of the transition metal to the lanthanide center could result in strong electronic and magnetic interactions.The concept of unsupported bonds between the lanthanide and the transition metal was first introduced and explored by Kempe et al.32. We demonstrate an important step toward molecular intermetallic nanomagnets—a perfectly trigonal planar [ErIII(ReICp2)3] complex  with magnetic hysteresis up to 7.2 K (at 2.2 mT s−1 magnetic sweep rate) where ErIII is solely coordinated by three diamagnetic ReI ions.Here, we apply the principles of easy-axis magnetic anisotropy design with the methodology of Kempe et al. that led previously to rare-earth metal complexes coordinated solely by transition metalsErRe3 is obtained following the literature procedure for [SmIII(ReICp2)3]32. The starting material [ErIII(btmsm)3] methyl anion) reacts with [Cp2ReH]49 in benzene methane  and the formation of the anionic [Cp2ReI]- species that readily coordinate to ErIII resulting in the formation of a trigonal planar complex [ErIII(ReICp2)3].ene Fig. . The btmErRe3 crystallizes slowly from benzene in the form of small yellow prism crystals which were characterized structurally by single-crystal X-ray diffraction 3] –3.056(13) Å range—well beyond typical coordination bonds of lanthanide complexes. The Er-Re distances in ErRe3 (2.9004(5), 2.9124(5), and 2.9172(5) Å) are similar to those reported for [SmIII(ReICp2)3], [LaIII(ReICp2)3], and [LuIII(ReICp2)3]32 as well as other molecular compounds with unsupported rare earth—transition metal bonds38 and the intermetallic SmCo5 with Sm-Co distances of 2.888 Å1. The comparison of the Re-Cp distances in ErRe3 and the [Cp2ReH] starting material  Å is similar to other compounds of this type52 and shows slight shortening of these contacts in ErRe3 to av. 1.864(8) Å 3] molecules are stacked on top of each other in a hexagonal fashion along the c crystallographic direction, leading to the formation of channels filled with disordered benzene molecules  Å. There are six more nearest ErIII neighbors within the 10 Å radius with distances larger than 9 Å 3] molecules that is sufficient for the observation of slow magnetization dynamics. Before attempting the magnetic characterization ErRe3 was subjected to a rigorous purity verification by performing a powder X-ray diffraction (PXRD) experiment for a sample loaded into a 0.3 mm glass capillary immersed in benzene and sealed using high vacuum grease (similar conditions were applied during the magnetic measurements discussed below). The experimental PXRD pattern shows very narrow peaks and matches almost perfectly the simulated one using 270 K SCXRD data 3] Fig.  with theule Fig. . There a(8)–3.0563 Å range(8)–3.0563 Å rangerIIIReICp3]  value of 11.3 cm3 K mol−1 at 260 K is close to 11.48 cm3 K mol−1 expected for a single ErIII 4I15/2 ground multiplet 3] (6H5∕2 multiplet)32. This is further confirmed by a very good agreement of the experimental χT and the ab initio calculations  dependence shows an abrupt increase below 18 K, which can be ascribed to ferromagnetic intermolecular interactions and blocking of magnetization.The unusual coordination environment of ErM(H)  measurements (magnetic field sweep rate of 2.2 mT s−1) revealed the presence of waist-restricted (pinched) magnetic hysteresis loops up to 7.2 K   s−1, C = 0.137(9) K−n s−1, n = 3.60(2), τ0 = 8.3(8)·10−10 s and Ueff = 285(3) K with R2 = 0.99994; Fig. HDC = 0.15 T could be reliably fitted including only the Raman and Orbach relaxation mechanisms using Eq. τ−1 is the magnetic relaxation rate, C and n parameters describe the Raman-like process, and τ0 and Ueff are related to the Orbach process. The satisfactory fit could be obtained with C = 0.0035(4) K−n s−1, n = 4.74(3), τ0 = 3.0(5)·10−10 s, and Ueff = 314(5) K with R2 = 0.99987  magnetic susceptibility measurements performed in the 1–10000 Hz frequency range up to 27 K in the absence of the external magnetic field  Fig.  containss A = 179 s−1, C =987 Fig. . The relUeff = 314(5) K is significantly smaller than expected from the theoretical calculations , predicting the fastest relaxation between the 5± Kramers Doublet lying 576 K above the highly axial ground state 2)3] characterized by Ueff = 122 K57 and surpasses the theoretical limit of 306 K (212.6 cm−1) for any trigonal planar molecule based on the theoretical analysis of a hypothetical ErIII(NH2)3 compound65. ErRe3 is a clear demonstration that direct bonding of transition metals to lanthanides might be one of the most promising directions towards room temperature MNMs, especially if paramagnetic donor atoms could be used.The obtained effective energy barrier for the magnetization reversal eff = 314 K is sigErRe3 are consistent with the recent data analysis by Gaita-Ariño et al12. The analysis shows relatively high magnetic hysteresis temperatures for prolate systems such as equatorial ErIII SIMs (coordinated by rigid ligands) accompanied by the maxima of the out-of-phase magnetic susceptibility (χ“) at relatively low temperatures. ErRe3 is a perfect demonstration of this trend with the Thyst of 7.2 K and Ueff = 314 K.Our findings regarding the magnetic dynamics of III(ReICp2)3] constitutes a molecular nanomagnet that mimics structurally the hardest commercial rare-earth magnet known SmCo5 and creates a connection between two incompatible fields: the rare-earth intermetallics and molecule-based compounds. ErRe3 is a rare example of a lanthanide single-molecule magnet coordinated solely by transition metals which enables the formulation of a promising strategy towards molecular intermetallic nanomagnets, where the lanthanide—the source of strong magnetic anisotropy—is directly coupled to the spins of softer transition metal ions playing the role of ligands.In conclusion, the reported triangular molecule [Er2 < 0.1 ppm and H2O < 0.5 ppm). All solvents used in this study (except benzene) were of HPLC grade and were additionally dried under argon using the Inert PureSolv EN7 solvent purification system and stored over 3 Å molecular sieves for at least 24 h before use. Anhydrous benzene was purchased from commercial sources (Merck) and degassed by performing three freeze-pump-thaw cycles before use. Anhydrous ReCl5 (at least 99.9%) was purchased from Merck and used as received for the preparation of [Cp2ReH] according to the modified literature procedure49 described below. The obtained [Cp2ReH] was characterized by means of SCXRD and the purity was confirmed by performing PXRD measurements. [ErIII(btmsm)3] was also prepared following the modified literature procedures48—see below for the description of the preparation procedure. The purity of the obtained product was checked by PXRD. See All manipulations were performed under an argon gas atmosphere inside the Inert PureLab HE glovebox (OIII(btmsm)3] was prepared following the modified literature procedures48. Inside the Ar-filled glovebox bis(trimethylsilyl)methyllithium66  was slowly added in small portions to a stirred clear pink solution of [ErIII(BHT)3]48  in 17 ml of pentane. After the addition is finished the reaction mixture is still slightly turbid due to the undissolved Libtmsm, but becomes clear after ca. 20 min of stirring and then turbid again after ca. 1 h of stirring due to the precipitation of the byproduct LiBHT. The stirring is continued for 5 h at room temperature and then the thick suspension is vacuum filtered through a P4 fritted glass funnel in order to remove LiBHT. Please note that the funnel needs to be thoroughly clean, rinsed with THF, and dried at 140 °C prior to use. The product might start to crystallize in the filtrate during vacuum filtration. The pink filtrate is transferred into a tightly sealed glass vial and kept at −40 °C for 5 h for crystallization. The obtained pink needle crystals are filtered by vacuum filtration using a PTFE membrane, collected, and stored in the freezer. Yield: 350 mg (39%). Additional crop of the product  is obtained by washing the LiBHT byproduct collected after the reaction with ca. 10 ml of pentane and keeping the washings at −40 °C for 24 h. The purity of both crops was confirmed by collecting PXRD patterns, which were identical to that simulated from the scXRD data.[Er2ReH] was synthesized according to the modified literature procedure49. Inside the Ar-filled glovebox, a 100 mL 29/32 Schlenk flask was charged with a large magnetic stirrer, sodium cyclopentadienide , and 25 mL THF, and was cooled to −40 °C. The cold mixture was set on a fast stirring and ReCl5  was added in small portions. The resulting mixture was left to heat to room temperature, then NaBH4  was added, the flask was sealed with a glass stopper, and was removed from the glovebox. Using standard Schlenk line techniques, the mixture was refluxed under Ar for 2.5 h, then it was cooled to room temperature and the solvent was removed under vacuum. The resulting brown solid was dried under vacuum for 30 min at 70 °C, then the reaction flask was fitted with a 29/32 cold finger filled with dry ice/acetone mixture and the brown solid was sublimed under vacuum. The temperature during the sublimation was gradually increased up to 180 °C under vacuum (p ≈ 10−2 mbar) and continued until the product was deposited on the cold finger surface as a brownish-yellow powder, leaving most of the brown residue on the bottom of the reaction flask. The reaction flask was moved inside the glovebox, the yellow product was moved to another Schlenk flask and the vacuum sublimation was repeated at 120 °C (this time with ice-cooling of the cold finger) yielding Cp2ReH as yellow needles. Yield: 0.50 g (26%). The purity of the product was confirmed by collecting the PXRD pattern, which was identical to that simulated from the scXRD data.[Cp2ReH]  was dissolved in anhydrous benzene (3.5 ml) and added to the benzene solution (2.5 ml) of [ErIII(btmsm)3] . The resulting orange solution was stirred using a glass rod for 5 min and then stored in a 20 ml scintillation vial that was left open for 5 h inside the glovebox. After this time yellow-orange crystals appeared which were collected by vacuum filtration using a 1-μm Teflon membrane (Whatman). Yield: 70 mg . The purity of the compound was checked by powder X-ray diffraction, with the experimental pattern : 796 s, 808 s, 816 s, 832 m, 850 m, 987 s, 1000 s, 1059 m, 1087 s, 1095vs, 1154vw, 1188vw, 1201vw, 1255 m, 1273 m, 1342 m, 1358 m, 1394 s, 1419 s, 1462 m, 1479 m, 1558 s, 1657s, 2001vw, 2034vw, 2854 s, 2929vs, 3036 m, 3063 s, 3089 s, 3550b . [CpSupplementary InformationPeer Review File


Herein, the intrinsic catalytic properties of unsupported In3Pt2, In2Pt, and In7Pt3 as model systems for Pt/In2O3‐based catalytic materials in MSR are addressed. In2Pt was identified as the essential compound responsible for the reported excellent CO2‐selectivity of 99.5 % at 300 °C in supported systems, showing a CO2‐selectivity above 99 % even at 400 °C. Additionally, the partial oxidation of In7Pt3 revealed that too much In2O3 is detrimental for the catalytic properties. The study highlights the crucial role of intermetallic In−Pt compounds in Pt/In2O3 materials with excellent CO2‐selectivity.Heterogeneous catalysts are often complex materials containing different compounds. While this can lead to highly beneficial interfaces, it is difficult to identify the role of single components. In methanol steam reforming (MSR), the interplay between intermetallic compounds, supporting oxides and redox reactions leads to highly active and CO Catalytic properties of In‐Pt intermetallic compounds in methanol steam reforming have been investigated. In2Pt has been proven to be mandatory for a high activity and selectivity. Only small amounts of In2O3 have been found to be beneficial for catalytic performance. It can be concluded that the decomposition of In7Pt3 into In2Pt and In2O3 leads to a less active and selective state in the form of an In2O3‐enriched material. However, for In2Pt and In3Pt2 surface sensitive analysis of sample composition is mandatory as decomposition of the near‐surface region cannot be identified by XRD.From the temperature‐dependent measurements, apparent activation energies of 2Pt sample reveals an asymmetric signal in the Pt4f region with a binding energy of roughly 71.8 eV 3.[2Pt. Similar results were obtained during the operando measurement. No changes in the Pt4f signal except for the removal of the underlying In4p signal are observed. In the In3d signal, a small shift to lower binding energy for the oxidic species is detectable, possibly resulting from the formation of partially reduced species under reaction conditions. In comparison to the UHV measurements, the relative amount of oxidic indium is decreasing, which confirms the stability of In2Pt under MSR conditions, as no continuous oxidation of the bulk occurred. Since investigations on supported In2Pt/In2O3 showed that oxidic indium is actively participating in MSR,2Pt alone but are the result of In2Pt, small amounts of oxidic indium and possibly a more Pt‐rich (inter)metallic surface species, since In2Pt is described as a line compound without significant homogeneity range.XPS measurements were conducted to ascribe the differences in the observed catalytic properties to the corresponding surface composition, with a focus on the formation of surface oxides. The as‐prepared InV Figure . The sig In(OH)3.42, 433Pt2 . Thus, it can be concluded that the observed catalytic properties of In3Pt2 are resulting from oxidic indium species and a Pt‐enriched surface species on top of bulk In3Pt2 and cannot be assigned to In3Pt2 alone, analogous to In2Pt.Almost identical results were obtained for In2 Figure . As for 2Pt and In3Pt2 show a slight surface oxidation in the as‐prepared state and under reaction conditions, it has to be concluded that the excellent CO2‐selectivity of In2Pt cannot be assigned to the formation of the surface oxides alone but depends strongly on the intermetallic compound. Besides the surface oxides, In2Pt is mandatory to obtain an excellent CO2‐selectivity, thus making In2Pt superior to In3Pt2. However, using In7Pt3 as precursor for In2Pt and In2O3 does not to lead to the excellent catalytic properties of bulk In2Pt with slight surface oxidation. From this, it can be concluded that the amount of In2O3 has to be limited on the In2Pt surface to obtain excellent catalytic properties. The pronounced differences between in the catalytic properties of In2Pt and In3Pt2 clearly show that the presence of In2Pt is an essential criterium for high CO2‐selectivity in catalytic In−Pt materials. In3Pt2, as oxidation product of In2Pt, is most likely also relevant for the high CO2‐selectivity but the obtained data clearly shows that In2O3 and In3Pt2 are not responsible for the excellent catalytic properties alone.Since both In3Pt2, In2Pt and In7Pt3, were synthesized as bulk materials and investigated regarding their catalytic properties and structural stability in methanol steam reforming. By operando TG/MS and XPS investigations, it was shown that In3Pt2 and In2Pt are stable under reaction conditions and only exhibit slight surface oxidation, whereas In7Pt3 decomposes into In2O3 and In2Pt. Upon linking these findings with supported In2Pt/In2O3,2‐selectivity of more than 99 % of In2Pt, which is significantly outperforming In3Pt2 and In7Pt3, can be ascribed to the presence of In2Pt and a small amount of oxidic indium. Large amounts of In2O3 are detrimental to the activity and selectivity of In2Pt, as observed for the strong decomposition in the case of In7Pt3. This study reveals that In2Pt, in combination with small amounts of In2O3 and In3Pt2 as decomposition products, is responsible for the high CO2‐selectivity of In−Pt materials in MSR and confirms the high capability of intermetallic bulk materials to understand the intrinsic roles of different compounds in heterogeneous catalysts.Three In‐rich intermetallic compounds, InThe authors declare no conflict of interest.1As 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.Supporting InformationClick here for additional data file.