First principle calculations show the prominent anharmonic coupling of the infrared inactive interlayer shear and the infrared active phonons. The nonlinear coupling terms generates an effective anharmonic force which drives the interlayer shear mode and lowers the ferroelectric switching barrier depending on the amplitude and polarization of infrared mode. Lattice dynamics simulations show that the interlayer shear mode can be coherently excited to the switching threshold by a train of infrared pulses polarized along the zigzag axis of MoS2. The results of this study indicate the possibility of ultrafast ferroelectricity in stacked two-dimensional materials from the control of stacking sequence.We theoretically investigate the mechanism of ferroelectric switching via interlayer shear in 3R MoS However, studies on the 2D ferroelectrics are yet an emerging field. Only a few materials such as SnTe5, α-In2Se36 have been experimentally shown to work in agreement with theories9. Transition metal dichalcogenide (TMDC) is another major 2D materials class showing versatile electronic phases ranging from semiconducting or metallic phase to that with topological characteristics12. Nonetheless, the ferroelectricity of the TMDC is very scarce due to the underlying symmetry of single layer in stable phase, which is either non-polar or centrosymmetric, precluding the electric polarization; i.e. 2H (P2/m) phases11. While theory showed ferroelectric instability of single layer 1T MoS213, its realization in the experiment is challenging because MoS2 is stable in semiconducting 2H phase rather than metallic 1T structure14.Ferroelectric two-dimensional (2D) materials are of great importance in realizing non-volatile devices with extreme feature sizeP3m1 for finite layers and R3m for bulk), hence the vertical electric polarization manifests itself in accordance with the global polar symmetry15. The direction of polarization depends on the stacking sequence, hence is reversed by the interlayer translation between the AB and AC stackings16 as shown in Fig. 2 showed switching of the polarization18 and the topological phase12 in the recent experiments, probably via the interlayer translation. The stability of the 3R structure MoS2 is comparable to that of the 2H structure and can be selectively synthesized among competing polytypes19. It is hence a viable candidate for 2D ferroelectrics in which the fascinating phenomena such as high electron mobility16 and valleytronics19 can be explored altogether.The ferroelectricity is likely to appear in stacked TMDC rather than in single layer form. It was recently shown that the horizontal mirror symmetry of individual layers is broken by the stacking in 3R structure  without the material damage32. The optical ferroelectric switching is a rapidly growing topic, which will enable the ultrafast and nondestructive way to achieve coherent switching35.The interlayer translation universally manifests as Raman active low-frequency lattice vibrations in layered 2D crystals due to the weak van der Waals bonds (vdW) between layers2 using the intense light pulse through the anharmonic phonon coupling. Density functional theory (DFT) calculations demonstrate that a large amplitude vibration of infrared mode can effectively lower the ferroelectric switching barrier, and induce an unidirectional anharmonic force on the interlayer shear mode along the switching direction. This effect depends on the polarization angle of the incident light pulse with respect to the crystallographic axis of MoS2 according to the selection rule. Lattice dynamics simulations indicate the possibility of dynamical ferroelectric switching through the coherent amplification of the interlayer shear mode and the lowering of the energy barrier under the repetitive pulses within a few picoseconds.In this work, we theoretically show the possibility of the ferroelectric switching of the bilayer 3R MoS2 has the polar point group symmetry of C3v with the polar axis along the z-axis. The 3R structure can be constructed by either AB or AC stacking sequence, which develops the spontaneous electric polarization in the opposite direction of each other as shown in Fig. P = 0.24 μCcm−2, in agreement with previous reports16. Figure NEB = 15.0 meV.The bilayer 3R MoS24. The bilayer 3R MoS2 has 18 zone-center phonon modes which are decomposed into Γ = 6A1 + 6E representations. The singly degenerate A1 mode involves out-of-plane motion of atoms, while the doubly degenerate E mode involves in-plane motion of atoms. Figure LS = 0.6 THz in phonon dispersion in Fig. QLS) involves the relative motion between adjacent layers along the in-plane polarization axis r. Therefore, the QLS mode is related to the AB ↔ AC stacking change. The infrared activity of the v-th mode is proportional to the square of mode effective charge 36. Due to the almost rigid relative ionic motion, the QLS mode does not produce net dipole moment as the calculated effective charge eμ−1/2 (where e is the electronic charge and μ is the atomic mass unit). The vanishing infrared activity indicates that it is almost impossible to directly excite the QLS mode with large amplitude to induce the stacking change. Nonetheless, the anharmonic coupling of QLS mode with other infrared active modes can provide an alternative route to control this mode, and the consequent ferroelectric switching. Among the other in-plane modes, only the high-frequency mode (denoted by QIR) at ΩIR = 11.4 THz shows finite effective charge eμ−1/2, and is the solely infrared active mode under vertical incidence of light.The optical switching mechanism was investigated based on the nonlinear phononicsQv of v-th mode is related to the atomic displacement vector mi is the atomic mass of i-th atom and QLS and QIR modes as {QLSx, QLSy} and {QIRx, QIRy}, respectively. They correspond to the linear polarization along the zigzag (x-axis) and armchair (y-axis) axes shown in the top view of the AB stacking of the 3R MoS2 in Fig. QLSy = +16.29 ÅQLSy = −16.29 ÅQLS along three crystallographically equivalent directions, r1 (+y direction), r2 and r3 directions (−120° and +120° from the +y direction), equally change the AB stacking into AC stacking as shown in Fig. The normal-mode coordinate QLS mode can be amplified along the desired direction for the AB ↔ AC stacking change by the coupling with QIR. Here, the anharmonic coupling property was investigated from the potential energy surfaces as a function of normal-mode coordinates. The potential energy surfaces V for each QIRx and QIRy were calculated using DFT on 21 × 21 × 23 points with steps of 0.82 ÅQIR and 1.63 ÅQLS modes. The energy surface was then fitted to the polynomial function asQIR is either QIRx or QIRy, and clmn is the anharmonic coefficient, and l, m and n powers of the normal-mode coordinates, respectively. Using this expression, we analyze the effect of the irradiating light pulse with the linear polarization along the x- or y-axis, thus exciting QIRx or QIRy mode, respectively. Note that the normal-modes in the cartesian basis are classified into the odd parity modes (QLSx and QIRx) and even parity modes (QLSy and QIRy) under the mirror symmetry σ1 shown in Fig. clmn is nonzero only for l + m = even (m = even). The terms were included up to 15th power terms (l + m + n = 15) in the polynomial function, which fits the DFT potential energy surface accurately. The representative coupling terms are displayed in Table The possibility of the ferroelectric switching hinges on how much the V represented on  coordinates when the amplitude of QIR is zero. The energy contour shows the directional dependence inherited from the C3v symmetry. The energy barriers for the AB → AC change along the equivalent r1,r2 and r3 directions in this potential energy surface are the same as Φ0 = 17.3 meV. The difference between Φ0 and ΦNEB for the AB → AC stacking change is because the deformation by the in-plane QLS mode does not include any out-of-plane relaxation, while the NEB path includes the relaxation from the slight increase (~1.6%) of the interlayer distance, reducing the barrier. It is worth to note that the Φ0 rather than the ΦNEB is relevant to the ultrafast switching in the picosecond time scale, while the latter is relevant to the conventional switching in a longer time scale.Figure QIR amplitude. Figure QIR was set to ±5.00 ÅC3v symmetry of the potential surface on the  coordinates was broken, and the energy barriers along the three equivalent directions became different. Under the negative amplitude of QIRx = −5.00 År1 direction decreased to 8.6 meV, but that along the other directions increased to 19.4 meV (r2 direction) and 21.3 meV (r3 direction), respectively. The energy landscape for the positive amplitude QIRx = +5.00 Åσ1. For both signs of QIRx, the coordinate of the potential energy minimum is slightly shifted along the r1 direction from the origin  direction. The energy landscape is symmetric with respect to the mirror plane σ1, and the energy barrier along the r2 and r3 directions are reduced . The dynamics of the nonlinearly coupled modes were simulated by the following coupled equations of motion,γIR and γLS are the damping coefficients for each mode, and F(t) is the optical driving force on the QIR mode. We used Gaussian pulse E0 is the amplitude of the electric field, σ is the duration of the pulse and Ω is the frequency.The dynamical behavior of the normal-modes was investigated under light pulse with a specific polarization direction. Since the motion of the QLS coordinate along the AB → AC switching directions when the QIR mode is resonantly pumped by a pulse with Ω = ΩIR, E0 = 34 MVcm−1 and σ = 100 fs. Such high intensity of the pulse is required to achieve the large amplitude of QIR  corresponds to the static case  direction . This explains the increase in the energy barrier along the r1 direction by 30% and the decreases along the r2 and r3 directions by 16%.The polarization-dependent modulation of the effective potential energy can be explained by the characteristics of the anharmonic coupling terms. The overall trend is captured by the coupling terms in the form of Fanh has a geometrical origin related to the Mo and S sublattices, which are displaced by the QIR in the opposite direction  in the adjacent layers as illustrated in Fig. k is spring constant and QIRx. This simple picture explains the Fanh along the +y direction, and agrees with the selection rule (l + m = even) for QIRy motion induces net forces along the −y direction due to the imbalance of the force (see the length of orange arrows).The polarization-dependent direction of ion Fig. . The motQIRx nor QIRy imparts such an unidirectional force on QLSx since the relevant coupling terms (the l) are absent due to the odd parity of QLSx. It prohibits the excitation of the interlayer shear along the r2 and r3 directions. Although the y-polarized light pulse lowers the energy barrier along the r2 and r3 directions, it cannot induce the ferroelectric switching along these directions. Therefore, the most effective way to realize the ferroelectric switching is to use the x-polarized light pulse which induces both the interlayer shear motion and energy barrier lowering along the r1 direction for the ferroelectric switching to occur.Neither QIRx-QLSy coupling under the x-polarized pulse. First, we considered a case neglecting the damping of normal modes to simply show the essential consequences of the QIRx-QLSy coupling on the dynamics of QLSy mode. Figure QIRx and QLSy modes at 0 K and 300 K, respectively, without damping under the x-polarized pulse with E0 = 34 MVcm−1 and σ = 100 fs. The initial vibration amplitudes were set as the mean-square-displacement QLS was assumed to be aligned to the y-axis by setting the initial coordinate as QLSx = 0 ÅQIRx = 0.21 ÅQLSy = 0.91 ÅQIRx = 0.28 ÅQLSy = 5.70 ÅNext, we analyze the dynamics of ferroelectric switching based on the QIRx and QLSy modes oscillate with the harmonic frequencies before the arrival of the pulse at 0 ps. The initial kinetic energy of QLSy mode was QIRx was pumped, QLSy started to oscillate with larger amplitude with respect to the shifted minimum at QLSy = 0.87 ÅQLSx . The eight sequencial pulses are applied to substantially amplify the QLSy mode from the zero-point vibration at 0 K. The time interval between the subsequent pulses is gradually increased by ~4% from the 1/ΩLS ~ 1.6 ps for the phase matching between the pulse and QLSy mode, considering the increase in the period of QLS mode due to anharmonicity. Upon each cycle of pulse irradiation, the QLSy mode is coherently amplified by gaining kinetic energy. After the eight pulses are irradiated, the QLSy mode has sufficient kinetic energy and jump over the effective barrier which is reduced by the QIRx mode. Once the initial AB stacking sequence changes to the AC stacking, it maintains the AC stacking due to the dissipation of the kinetic energies of the vibrations. This corresponds to the AB → AC ferroelectric switching. The opposite switching operation, AC → AB, can be performed by the same optical input as illustrated in Fig. Fanh on QLSy mode in the AC stacking is reversed (−r1 direction) with respect to that (r1 direction) in the AB stacking. The optical parameters of pulses (e.g. E0 and σ) used in this study might be optimized further for more efficient switching, for instance, via pulse shaping techniques37.Secondly, a more realistic model that includes the damping of the normal modes was considered. The damping coefficients of 2 whose direction of the polarization is reversed by the change of the stacking sequence was investigated. The ferroelectric switching was achieved by driving the interlayer shear mode through the anharmonic energy flow from the optically pumped infrared mode. Remarkably, due to the selection rule from the crystal symmetry of MoS2, the degenerate interlayer shear mode can only be driven along its armchair axis whether the infrared mode is pumped along the zigzag or armchair axis. However, the optical pulse should be polarized along the zigzag axis for successful switching since the direction of anharmonic force is aligned with the switching direction. The coherent light pulses can amplify the interlayer shear mode substantially and unidirectionally, displacing the stacking sequence into the opposite polarization. The scheme for optical modulation of the stacking structure can be applied to other 2D materials exhibiting the interlayer shear mode to explore various stacking-dependent properties in a dynamical manner.In summary, the polarization switching mechanism of the bilayer 3R MoS39. The projector-augmented wave (PAW) method40 and a cut-off energy of 500 eV were used with the valence electron configurations of Mo[4s24p65s24d4] and S[3s23p4], respectively. The generalized gradient approximation41 with Grimme’s D3 scheme42 was used to describe the van der Waals interaction. The bilayer structure was simulated by the supercell containing ~40 Å of vacuum layer to avoid the artificial interaction between periodic images. The structures were fully relaxed using 24 × 24 × 1 k-mesh until the residual forces on the atoms were less than 0.001 eVÅ−1. Spontaneous polarization was calculated using Berry phase method43. The phonon calculation was performed using PHONOPY code44 using 3 × 3 × 1 supercell and 8 × 8 × 1 k-mesh.The density functional theory calculations were performed using Vienna Ab-initio Simulation Package (VASP)


Incorrect versions of S1–S3 Tables were published in error. This article was republished on September 27, 2019 to correct for these errors. Please download this article again to view the correct version.


J–V ) hysteresis. Inorganic perovskite nanocrystals (IPNCs) are promising candidates for high‐performance photovoltaic devices due to their simple synthesis methods, tunable bandgap, and efficient photon downshifting effect for ultraviolet (UV) light blocking and conversion. In this work, CsPbBr3 IPNCs modification could give rise to the vapor phase and solution‐processed PSCs with a power conversion efficiency (PCE) of 16.4% and 20.8%, respectively, increased by 11.6% and 5.6% compared to the control devices for more efficient UV utilization and carrier recombination suppression. As far as is known, 11.6% is the most effective enhanced factor for PSCs based on photon downshifting effect inside of devices. The CsPbBr3 layer could also significantly retard light‐induced degradation, leading to the lifetime over 100 h under UV illumination for PSCs. Additionally, the modified PSCs exhibit weak hysteresis and multiple colors of fluorescence. These results shed light on the future design of combining a photon downshifting layer and carrier interfacial modification layer in the applications of perovskite optoelectronic devices.Photovoltaic devices employing lead halide perovskites as the photoactive layer have attracted enormous attention due to their commercialization potential. Yet, there exists challenges on the way to the practical use of perovskite solar cells (PSCs), such as light stability and current–voltage ( Their superior quality in luminescence quantum yield, size‐tunable bandgap, high transparency in visible spectrum, and excellent light stability makes the CsPbX3 IPNCs attractive and promising in luminescent downshifting application in PSCs. When sunlight illuminates on the IPNCs, the UV light can be converted into visible light by the down‐conversion effect. Then the converted visible light can be absorbed efficiently by perovskite photoactive films. Thus, the IPNCs are deemed to be helpful in enhancing the PCE and light stability of PSCs. In this work, IPNCs were prepared using a hot‐injection method and then successfully applied to the PSCs inside and outside. Herein, we construct an appropriate IPNCs interlayer via spin‐coating method, which acts as a seed‐mediated layer for perovskite film formation, with both uniform morphology and favorable absorption capacity. In addition, IPNCs and hybrid perovskites film form a fully crystalline heterojunction, which is beneficial to minimizing the defect and trap densities. The electrical mechanism of assembled nanocrystals inside the devices was also investigated. The results reported here reveal that making full use of the UV light by CsPbX3 nanocrystals is beneficial for enhancing the photovoltaic performance of solar cells.An ideal luminescent downshifting layer needs to have strong UV light absorption, excellent luminescence quantum yield, extraordinary transparency in the visible region, and large Stokes shift to reduce the self‐reabsorption losses. Moreover, structural simplicity, low‐cost fabrication, ambient operation, and long‐term stability are also crucial for real applications. Commonly, lanthanide‐based inorganic materials  images in Figure S1 in the Supporting Information show cubic morphology of as‐prepared CsPbBr3 nanocrystals with an average size of 8 nm. Then, IPNCs with narrow size distribution were redispersed and purified in ethyl acetate before being deposited on the FTO/cp‐TiO2 substrate. In the purification process, the surface ligands such as oleic acid and ammonia groups can be eliminated by ethyl acetate. These ligands block the transport of photo‐generated carrier, leading to enhanced electron mobility and improved device performance.Figure3 in ethyl acetate under UV light irradiation. The purified nanocrystals can still retain the cubic morphology for several months in ambient conditions  spectra in Figure 3 film by ethyl acetate with different purification time.3 was introduced into the devices. One to seven layers of purified CsPbBr3 were prepared on the FTO substrate by spin‐coating method to optimize the performance. Then the CH3NH3PbI3 films were directly deposited on top of the precoated CsPbBr3 layer using a solvent‐free, dual‐source vapor deposition method to avoid the dissolution of IPNCs in solvent , dimethyl sulfoxide (DMSO) et.). Simultaneously, dual‐source vapor deposition method for CH3NH3PbI3 lay preparation could avoid annealing process, which is necessary for solution method.3‐modified CH3NH3PbI3 perovskite thin films were collected in Figure 3 films, the XRD patterns exhibit typical peaks of orthorhombic CsPbBr3 around 15° and 30°. It also shows no phase transition in the IPNCs structure during ethyl acetate purification. In order to detect changes inside the perovskite film more quickly, the thickness of CH3NH3PbI3 was reduced from 470 to 150 nm. XRD patterns were performed immediately after CH3NH3PbI3 thin film was vapor deposited. After 45 min of depositing 150 nm CH3NH3PbI3 film, the XRD patterns of the film exhibits a relatively strong CsPbI3 peak in Figure 3 peak, indicating that most of the Br ions in CsPbBr3 has been replaced by I ions from CH3NH3PbI3. Further depositing 470 nm CH3NH3PbI3 films on FTO/CsPbBr3 substrate for one day, there are no obvious CsPbBr3 peaks. This can be attributed to the thicknesses of FTO (570 nm) and CH3NH3PbI3 layer (470 nm) which are thicker than that of CsPbBr3 IPNCs (110 nm). Furthermore, the CsPbBr3 IPNCs are surrounded by CH3NH3PbI3 films, so the peaks of CsPbBr3 are hard to be detected. More importantly, most of CsPbBr3 has been converted into CsPbIxBr3−x, and further into MAyCs1−yPbIxBr3−x alloy, through the ion exchange process. The diffraction peaks of FTO/IPNCs/CH3NH3PbI3 film at 2θ = 14.2°, 28.5°, and 32.0° are assigned to (100), (200), and (210) planes of tetragonal CH3NH3PbI3, respectively.2/IPNCs substrate, which is beneficial to improving interface contact and enhancing the PCEs.The optical properties of CsPbBrFigure2/IPNCs/CH3NH3PbI3/spiro‐OMeTAD/Au for the CsPbBr3‐modified PSCs. The morphologies of the deposited IPNCs on FTO/cp‐TiO2 substrate were studied by scanning electronic microscopy (SEM) in Figure 2 substrate. Further increasing the amount of IPNCs layers, more nanocrystals agglomerate and the surface roughness increases obviously, as illustrated in the atomic force microscopy (AFM) images . The root‐mean‐square (RMS) values from AFM image were calculated to be 7.8, 10.2, and 14.1 nm, revealing the enhanced surface roughness by increasing the layers. In comparison with the smooth surface of the unmodified TiO2 layer, the rough surface of the cp‐TiO2/CsPbBr3 layer could provide larger surface area and better contact with the hybrid perovskite layer, resulting in more effective channels for the carrier separation and collection. To give further insight in the structure of cp‐TiO2/CsPbBr3, the cross‐section images indicate that the thickness of IPNCs layer on the cp‐TiO2 was about 110 nm , corresponding to five times spin‐coating of CsPbBr3 solution. The surface morphology of CH3NH3PbI3 thin film is shown in Figure 3 IPNCs modification under UV light irradiation. S0 is a control device, while the S1 is the IPNCs‐modified device. Green color fluorescence is shown for the device with IPNCs only but without CH3NH3PbI3 layer. Red fluorescence is observed for the device with CH3NH3PbI3 layer deposited on cp‐TiO2/ CsPbBr3 substrate. This can be attributed to the ions exchange of I− in CH3NH3PbI3 with the Br− in CsPbBr3, resulting in the formation of CsPbBr3−xIx IPNCs and CH3NH3PbI3−xBrx, causing the red shift of the luminescence spectrum of CsPbBr3−xIx IPNCs and slight blue shift of the absorption spectrum of the CH3NH3PbI3−xBrx layer.3NH3PbI3 films on cp‐TiO2/CsPbBr3 substrate by the AFM images . Figure 2, CsPbBr3, and CH3NH3PbI3, respectively. In comparison with the unmodified devices, the CsPbBr3 modification could enhance the light absorption of hybrid CH3NH3PbI3 films.According to previous study, the uniformly distributed nanocrystals allow for proper control of perovskite film growth in terms of nucleation rate and new phase formation.3 modification is beneficial to the device performance, the devices with structure of FTO/TiO2/CsPbBr3/CH3NH3PbI3/Spiro‐OMeTAD/Au were fabricated and characterized. The current–voltage (J−V) curves of the PSCs with different layers of CsPbBr3 are presented in Figure3 layer may provide higher contact area for CH3NH3PbI3, the excessive CsPbBr3 will agglomerate on TiO2 electron transport layer, affecting the crystallization of hybrid perovskite layer and leading to the decreased device performance. This can be confirmed from Figure S6 in the Supporting Information, which reflects the influence of CsPbBr3 with different thicknesses on the absorption spectra of CH3NH3PbI3 thin film. When the spin‐coating times are less than four, the absorption intensity of CH3NH3PbI3 films increases in the full spectrum range, which indicates the improved crystallization. In addition, significant enhancement in the range of 200–500 nm can be observed, which is attributed to the enhanced absorption of IPNCs in blue and UV region. As the IPNCs spin‐coating times is more than five, the overall absorption of the device decreases, which can be explained by the poor crystallinity of hybrid perovskite films. As a result, an optimized thickness of CsPbBr3 layer is about ≈110 nm, leading to the best performance with a maximum PCE of 16.4% in this work, with enhanced factor of 11.6% compared to control device. To the best of our knowledge, this is the most efficient enhanced factor in PSCs compared to previous works, which is applied to the downshifting layer inside of PSCs.J–V curves of the optimized devices fabricated with and without CsPbBr3 are shown in Figure Jsc) of 23.31 mA cm−2 and fill factor (FF) of 68.9%, compared to those of control device (22.49 mA cm−2 and 61.0%). These results indicate that the CsPbBr3 layer has a significant effect on scJ, FF, and PCE. The open‐circuit voltage (Voc) of the CsPbBr3‐modified device is 1.02 V, which is lower than that of control devices (1.07 V). This can be related to changes of the bandgap in CH3NH3PbI3 after Br− (CsPbBr3) passivation effect.3 IPNCs), and the photovoltaic parameters are summarized in Figure Jsc, and FF compared to those of the control devices can be attributed to lower series resistance. The narrow PCE distribution of the modified devices in Figure S7 in the Supporting Information further indicates that modified devices present excellent reproducibility compared to control devices.In order to study whether the CsPbBr−1), as shown in Figure3 IPNCs could achieve perfect interfacial contact, facilitating the photo‐generated electron transfer by creating electrostatic interaction between TiO2 and CH3NH3PbI3.Previous results indicate the hysteresis effect of PSCs is clearly sensitive to interface between perovskite and electron transport layer.Rrec) was studied with impedance spectroscopy (IS) measurement by fitting the results using an equivalent circuit model of Rs(RrecCrec) in Figure S8 in the Supporting Information. Accordingly, the data at middle frequency in IS patterns are related to charge recombination processes. A higher Rrec indicates a lower recombination rate, and could result in lower carrier recombination in PSC devices.Rrec values obtained with different applied voltages are shown in Figure S9 in the Supporting Information and summarized in Figure Rrec, while the CsPbBr3‐modified devices exhibit higher Rrec at 0–0.5 V bias. By introducing IPNCs in PSCs, higher Rrec values indicate the enhanced carrier transient behaviors due to more efficient pathways for electron transport from CH3NH3PbI3 to TiO2 induced by the purified CsPbBr3, which causes the elimination of hysteresis effect and enhancement of FF.In order to further verify the effect of interfacial engineering on carrier transmission and hysteresis effect. The recombination resistances (2/CsPbBr3/CH3NH3PbI3/PCBM/Ag) to measure the electron trap‐state density to further study the enhanced transient carrier behaviors. Figure VTFL) using the equationnt is the trap‐state density, e is the elementary charge, L is the thickness of CH3NH3PbI3 (450 nm), ε is the relative dielectric constant of CH3NH3PbI3 (ε = 32),0 is the vacuum permittivity. The VTFL values of the control‐ and IPNCs‐based devices are 0.50 and 0.48 V, respectively. Therefore, the electron trap‐state density can be reduced by IPNCs modification.3‐modified CH3NH3PbI3 films exhibit higher conductivities than that of control films. At space‐charge limited current (SCLC) region, the modified device exhibits higher electron mobility. The obtained results are 3.07 × 10−4 and 2.36 × 10−3 S m−1 for the modified and control devices, respectively, which are comparable to the previous reports.3 IPNCs can benefit carrier transport between TiO2 electron transport layer and CH3NH3PbI3.We then fabricated electron‐only transporting devices  element analysis of cross‐sectional CsPbBr3/CH3NH3PbI3 films  was carried out and is depicted in Figure S11 in the Supporting Information. For the initial cross‐section EDX mapping, it can be seen that the main distribution thickness of Cs element is about 100 nm, with Br element distributed around Cs element. This indicates that the exchange of anion Br is faster than that of cationic Cs. After two hours, the EDX mapping from the same cross‐sectional SEM interface shows that the distribution of Cs element is widened to about 145 nm, indicating that Cs ions not only existed in CsPbBr3, but also entered into CH3NH3PbI3. At the same time, the distribution range of Br element further increased to some extent. It shows that some Br ions also enter into CH3NH3PbI3. The EDX mapping results indicate the ion exchange does occur between inorganic CsPbBr3 and hybrid CH3NH3PbI3 perovskite, leading to Cs1−yMAyPbI3−xBrx alloy between CH3NH3PbI3 and CsPbBr3 by the exchange mechanism of anions and cations. The possible mechanism diagram is shown in Figure 3 can effectively promote electron injection from CH3NH3PbI3 to cp‐TiO2 and reduce the recombination and quenching at the interface of CH3NH3PbI3/TiO2. According to previous reports, a small amount of Cs and Br can promote the crystallization of perovskite, leading eventually to improved optical and electrical properties of complex perovskite photoactive material.J–V hysteresis for the fewer capacitance elements in the traps.3NH3PbI3 to TiO2 layer can be enhanced, while the hysteresis effect can be reduced after CsPbBr3 IPNCs modification due to the decreased trap states and increased electron conductivity.Above all, electron trap‐state density measurement further confirms that the high electron extraction efficiency originated from improved photoelectrical properties induced by IPNCs modification. As an efficient scaffold layer, the CsPbBrFigure3 IPNCs‐modified device in the whole region from 400 to 800 nm in comparison with that of the control device. Notably, the modified devices show higher intensity in the blue and UV region (300−500 nm), which is ascribed to the increased photo response contributed by CsPbBr3 nanocrystals by photon downshifting effect.−2, close to the value from J–V curve. Thus, the excellent performance with reduced hysteresis achieved in the modified PSCs is attributed to the efficient carrier transport in comparison with that in the control PSCs.3 and CH3NH3PbI3 and transformation from CsPbBr3 nanocrystals to MAyCs1−yPbBr3−xIx films. We also tested the influence of CsPbBr3 layer on device stability under continuous 365 nm UV light illumination for 100 h in nitrogen environment. In this UV irridiation aging test, the unmodified device loses 45% of its initial PCE, dropping from 14.7% to 8.2% in Figure 3‐modified device demonstrates excellent stability under the same condition, retaining 82% of its initial PCE. This indicates that the device fabricated with IPNCs demonstrates excellent light stability. We also fabricated the device by coating IPNCs outside of the device by modify CsPbBr3 using silane surface coating method. Figure S12 in the Supporting Information shows the TEM image of CsPbBr3@SiO2 nanocomposite. According to previous work, silane surface coating can effectively improve environmental stability of IPNCs (CsPbBr3@SiO2 nanocomposites).3@SiO2 nanocomposites were measured to be 79.4%, which is slightly lower than that of pure CsPbBr3 nanocrystals (82.8%). The air‐stable CsPbBr3@SiO2 layer was coated on the front side (transparent side) of the device as shown in the schematic diagram for estimating the influence of UV blocking for the devices, which is mainly attributed to the carrier barrier effect of dielectric silicon dioxide on the surface of IPNCs.To study the photocurrent stability, the photocurrent densities of the champion devices were recorded as a function of time when the cells were biased at their respective maximum power point values, as shown in Figure 3NH3PbI3 precursor method. From the J–V measurements of the PSCs‐coated outside by IPNCs in comparison with control device in Figure Voc and FF values are almost the same, 1.06 V and 79.7%/79.6%, indicating that the effect of IPNCs coating on Voc and FF values of the devices is negligible. The Jsc exhibits an increase from 23.3 to 24.6 mA cm−2, as the optimized spin‐coating times is six. The enhanced Jsc results from the enhanced absorption by the photon downshifting effect of UV light to visible light by CsPbBr3@SiO2 nanocomposite layer. The PCE exhibits the same tendency as Jsc, and increases from 19.7% to 20.8%, with the amplification of 5.6% compared to the bare device, indicating that the enhanced performance from photon downshifting effect is about 5.6%. The external quantum efficiency (EQE) spectra from 360 to850 nm further indicate that CsPbBr3@SiO2 nanocomposite–coated devices could exhibit the improved light harvesting effects than that of the bare device, especially in the region of UV and blue light as depicted in Figure S15 in the Supporting Information. The integrated photocurrent densities are 22.08 and 23.0 mA cm−2 , with a 4.2% enhancement. The mismatched value inferred from the J–V curves  of its initial PCE after 1000 h irradiation under simulated light, while the bare device exhibits a dramatically decreased PCE to 64%. This stability result further proves the photon downshifting effect of CsPbBr3 and its blocking effect on UV light, which is beneficial to improve the light stability of the device. Here, it can be concluded that the CsPbBr3 between TiO2 and CH3NH3PbI3 play three important roles. First, at the interface between CH3NH3PbI3 and CsPbBr3 layers, the electron transfer process from CH3NH3PbI3 to TiO2 electron transport layer is improved, leading to increased PCE. Second, light harvesting in UV region is enhanced via photon downshifting effect of IPNCs to improved PCE. Third, IPNCs as a UV filter can significantly retard the UV photo‐induced degradation of hybrid perovskite films to increase light stability.To study the effect of IPNCs coating on the light stability under long time irradiation, the device with an excellent PCE of 19.7% is prepared by a typical one‐step spin‐coating CH33 IPNCs to improve the photovoltaic performance and light stability of PSCs by enhanced carrier transport and photon downshifting effect. Incorporating IPNCs layer (≈110 nm) leads to the device with a maximum PCE of ≈16.4%, with the amplification of 11.6% compared to the control devices. To the best of our knowledge, this is the first report about improving the PCE and stability of PSCs with the introduction of CsPbX3 nanocrystals inside the devices. This strategy could exhibit the highest enhanced factor of 11.6% compared with the traditional photon downshifting material. Additionally, the outside CsPbBr3@SiO2 coating could increase the PCE from 19.7% to 20.8%, which is improved by 5.6% compared to the uncoated device. More importantly, CsPbBr3 IPNCs‐modified PSCs exhibit much better stability under UV light irradiation compared to the control devices. The fundamental mechanism in optical and electrical properties reveals that the modification of IPNCs can provide noticeable advantages, including enhanced light harvesting and the electron injection from CH3NH3PbI3 hybrid perovskite layer to TiO2 electron transport layer. Furthermore, integrating IPNCs inside could induce specific color fluorescence and colorful perovskite devices under UV illumination. The IPNCs modification developed in this work exhibits significant potential toward the fabrication of PSCs with high PCEs.In this work, we present a study of the application of CsPbBr4Materials and Precursor Preparation: Chemicals for preparation like Cesium bromide , Cesium iodide , lead diiodide , lead dibromide , and Cesium carbonate  were obtained from the Alfa Aesar. Hydroiodic Acid  and Hydrobromic acid  were obtained from Aladdin Reagents Co., Ltd. . Oleylamine , and Octa‐decene  were purchased from MACKLIN Co., Ltd. (China). Spiro‐OMeTAD powders were purchased from Xi'an Polymer Light Technology Corp (China). Spiro‐OMeTAD hole transporting solutions were made by dissolving 72.3 mg of Spiro‐OMeTAD powders  in 1 mL of chlorobenzene (CB) and stirring it at 70 °C for 2 h. Then, 17.5 µL of Li‐bis(trifluoromethanesulfonyl) imide (Li‐TFSI) (520 mg mL−1 in acetonitrile) and 28.3 µL of 4‐tert‐butylpyridine (t‐BP) were added to the solution in a nitrogen‐filled glove box and stirred for 30 min until used for coating. The other materials and chemicals were purchased from MACKLIN Co., Ltd (China) and all chemicals were used as received. The CsPbBr3 IPNCs was prepared according to the previous reports. Specifically, the obtained precipitates were first washed with toluene through centrifugation at 9800 rpm for 10 min. Then the precipitates were further washed with ethyl acetate through centrifugation at 9000 rpm for 1 min. This washing process was repeated for 4 times to eliminate organic ligands on the surface as much as possible. Finally, the precipitates were redispersed in ethyl acetate and formed a stable colloidal solution.Device Fabrication and Characterization: All device fabrication was carried out in glove. FTO/glass substrate was etched with 30% HCl and zinc powder was removed for anode contact. Then the substrates were cleaned with detergent , deionized water, ethanol, and acetone. Finally they were treated under oxygen plasma for 15 min to remove organic residues. Cp‐TiO2 layers were deposited by spin‐coating of titanium (IV) isopropoxide (3.24 g) in butanol (50 mL) solution with a speed of 2000 rpm. After that, the substrates with titanium (IV) isopropoxide were annealed in air at 500 °C for 30 min. PbI2 and CH3NH3I powders were used to fabricate CH3NH3PbI3 films via thermal evaporation process. The evaporation current for PbI2 was 32 A by the inorganic vapor source with evaporation rate of 0.24 A s−1, and CH3NH3I was evaporated by organic vapor source at 320 °C with evaporation rate of 0.31 A s−1. The chamber pressure was controlled to be ≈6.4 × 10−4 Pa. The film thickness was detected by the film‐thickness meter in the chamber and controlled by the evaporation time. For the one‐step perovskite base devices preparation, PbI2 (462 mg) and CH3NH3I (163 mg) were first dissolved in mixed solvent  and stirred at 60 °C for 4 h before use. Then 100 µL mixture was deposited onto the FTO/cp‐TiO2 by a successive spin‐coated process of two steps, i.e 6000 rpm for 9 s and 4000 rpm for 40 s to form CH3NH3PbI3 film. Specifically, 600 mL diethyl ether was dropped on the surface. The second spin‐coating process was for 15 s. Then the antisolvent‐treated CH3NH3PbI3 films were dried at 100 °C for 10 min.J–V curves of the PSCs were obtained by a solar simulator (ABET Sun 2000) at AM1.5 and 100 mW cm−2 illumination calibrated with a reference silicon cell (RERA Solutions RR‐1002) using a Keithley 2400 as a source‐meter. IPCE was recorded by using SolarCellScan100 as a function of wavelength from 380 to 800 nm. The spectral EQE responses were taken from 360 to 850 nm by an EQE measurement system . The surface morphology of the films was characterized by a Sirion field‐emission scanning electron microscope. XRD measurements for the prepared films was carried out on a Rigaku D/max 2550 X‐ray diffractometer using a monochromatized Cu target radiation source at a scanning rate of 18° min−1. The electrical conductivity was measured by a Keithley 2400 as a source‐meter, with current–voltage traces from +3 V to −3 V.The solar cell measurement and storage were performed in ambient conditions, without any encapsulation. The The authors declare no conflict of interest.SupplementaryClick here for additional data file.


New battery chemistries based on LiOH, rather than Li2O2, have been recently reported in systems with added water, one using a soluble additive LiI and the other using solid Ru catalysts. Here, the focus is on the mechanism of Ru‐catalyzed LiOH chemistry. Using nuclear magnetic resonance, operando electrochemical pressure measurements, and mass spectrometry, it is shown that on discharging LiOH forms via a 4 e− oxygen reduction reaction, the H in LiOH coming solely from added H2O and the O from both O2 and H2O. On charging, quantitative LiOH oxidation occurs at 3.1 V, with O being trapped in a form of dimethyl sulfone in the electrolyte. Compared to Li2O2, LiOH formation over Ru incurs few side reactions, a critical advantage for developing a long‐lived battery. An optimized metal‐catalyst–electrolyte couple needs to be sought that aids LiOH oxidation and is stable towards attack by hydroxyl radicals.Non‐aqueous Li–O It was also shown on charging that the LiOH can be removed with the aid of LiI3 at around 3.1 V.2O2 and LiOH was formed on discharge, and that on charging, Li2O2 was first converted to LiOH, the latter then being decomposed by Ru catalysts at voltages of as low as about 3.2 V. At higher water contents (ca. 250 ppm), LiOH formation appeared to be dominant on discharge.2 battery, but fundamental insight into LiOH based chemistries may also aid in the development of Li2O2‐based batteries that operate utilizing air (or moist oxygen), where LiOH inevitably forms.Recently, LiOH has been identified as the major discharge product in a few of Li–O2 is involved in LiOH formation, this process incurring fewer side reactions compared to Li2O2. On charging, the LiOH is quantitatively removed at 3.1 V, with the oxygen being trapped in the form of soluble dimethyl sulfone in the electrolyte.Herein we develop a mechanistic understanding of the Ru‐catalyzed oxygen chemistry. Using quantitative nuclear magnetic resonance and operando electrochemical pressure and mass spectrometry measurements, we show that on discharging, a total of 4 electrons per O2 batteries prepared using Ru/SP electrodes with various concentrations of added water in a 1 m LiTFSI/DMSO (lithium bis(trifluoromethane) sulfonimide in dimethyl sulfoxide) electrolyte. In the nominally anhydrous case, discharge and charge plateaus are observed at 2.5 and 3.5 V, respectively, where an electrochemical process involving two‐electrons per oxygen molecule and Li2O2 formation dominates on discharging . As the water content increases, it is clear  carbon electrode is described in the Supporting Information. Microscopy and diffraction experiments show that Ru crystals of less than 5 nm are well dispersed on the SP carbon substrate . Figure 2 reduction rather than from electrolyte decomposition, we performed NMR experiments with isotopically labeled H (D2O) and O  . Up to four electrons can be stored per O2 molecule, the theoretical capacity of the battery operating via Reaction (1) being 1117 mAh/gLiOH, comparable to Li2O2 (1168 mAh/gLi2O2). To examine the role of Ru in LiOH formation further, we discharged a SP electrode in a 1 m LiTFSI/DMSO electrolyte with 4000 ppm water. XRD and SEM show that the discharge leads to mainly Li2O2 formation with an e−/O2 ratio of 2.2 , whereas discharging Ru/SP in the same electrolyte forms only LiOH. This contrasting behavior suggests that in the absence of Ru, the reaction between H2O and Li2O2, 2 Li2O2+2 H2O→4 LiOH+O2 (2), is slow, even though it is thermodynamically favorable (ΔG°=−149.3 kJ mol−1) but Ru clearly promoted the LiOH formation. By exposing a Ru/SP electrode discharged in a nominally dry electrolyte (where Li2O2 is the main product) to the 4000 ppm water‐added electrolyte, XRD  shows that all the Li2O2 was converted into LiOH in the presence of Ru after 10 h ; this indicates that Ru can catalyze reaction (2) given above. It is likely that the electrochemical formation of LiOH in the Ru/SP system proceeds via first Li2O2 generation (O2+2 e−+2 Li+→Li2O2) and then Ru catalyzes the chemical reaction of Li2O2 with H2O to eventually form LiOH (Reaction 2); overall the reaction is O2+4 e−+4 Li++2 H2O→4 LiOH. This observation also suggests that water present must be important to solubilize Li2O2, LiOH, and derived species and facilitate the solid‐solid phase conversion (from Li2O2 to LiOH).Operando pressure measurements show that the recorded pressure matches well with the trend line expected for 4 e1H solid‐state NMR spectra  clearly generated Li formate, acetate, methoxide side reaction products (signified by the resonances at 0–10 ppm),1H and 13C solution NMR measurements of electrolytes extracted after electrochemical discharge and from a slurry of LiOH and electrolyte held under O2 for 30 days also show that hardly any soluble side‐reaction products are formed .Importantly, the LiOH formation during discharge involves few parasitic reactions. Quantitative a Figure  E compar−2 (0.5 mAh or 1250 mAh/gRu+C per cycle), with consistent electrochemical profiles . Although the ex situ tests supported a highly reversible O2 electrochemistry, operando electrochemical pressure and mass spectrometry experiments suggested otherwise: very little gas was evolved on charging  was used in the electrolyte. Figure 1H (A), 13C (B), and 17O (C), and 1H–13C heteronuclear single quantum correlation (D) solution NMR spectra of the cycled electrolytes. New peaks at 2.99 ppm (1H), 42.6 ppm (13C), and 169 ppm (17O) appeared and progressively intensified with increased cycle number; these resonances consistently point towards the formation of dimethyl sulfone (DMSO2), its identity being further corroborated in the 1H–13C correlation spectrum. Of note, the 17O signal of DMSO2 is even stronger than the large amount of natural abundance (NA) DMSO used in the solution NMR experiment, suggesting that DMSO2 is likely to be 17O‐enriched. Its growth in intensity is accompanied by the decrease of H217O, indicating that some 17O from H217O ended up in DMSO2 owing to isotope scrambling in the charging process. Given that LiOH is quantitatively formed and then removed on charge ; DMSO+2 .OH →DMSO2+H2O (4). The overall reaction thus is: 2 DMSO+4 LiOH→2 DMSO2+2 H2O+4 e−+4 Li+ (5). It is seen that the same number of electrons is involved in discharge (Reaction 1) and charge (Reaction 5), with one O reacting per two electrons . The electrochemical process, Reaction (3), sets the voltage observed on charging, rather than the overall Reaction (5). Surface adsorbed hydroxyl species are generally considered to be the first reaction intermediates on many OER metal catalysts in aqueous media.2 and thus the battery can be continuously charged at a low voltage until all solid LiOH products are removed . The resulting DMSO2 is soluble in the DMSO electrolyte and will not immediately impede ion diffusion or interfacial electron transfer as other insoluble by‐products would do, which is perhaps why this side reaction does not rapidly lead to battery failure.Further solution‐NMR measurements were performed on electrolyte samples prepared from several charged cells extracted following different cycle numbers, where 2O2 to LiOH formation, similar reactions being seen for several other metal catalysts. The cell discharge reaction consumes four electrons per reduced O2 molecule. This LiOH formation process involves very few side reactions and LiOH itself is much more stable in organic solvents than Li2O2; these are the fundamental prerequisite for a long‐lived Li–O2 battery. On charging, the Ru quantitatively catalyzes LiOH removal via DMSO2 formation rather than O2 evolution. We propose that DMSO2 forms by the reaction of hydroxyl radicals with DMSO, the former being generated on Ru catalyst surfaces. This work highlights the advantage of using metal catalysts to catalyze a 4 e− ORR with very few side reactions, and also the unique role of a metal catalyst in promoting LiOH formation versus electrolyte decomposition. An optimized catalyst–electrolyte couple needs to be sought for to satisfy both activity towards LiOH oxidation and stability against electrolyte decomposition on charging. This work provides a series of key mechanistic insights into the Ru‐catalyzed Li–O2 battery in the presence of water, which will aid the design of catalyst and electrolyte systems that can be used in more practical batteries.In summary, we have shown that with added water (beyond 4000 ppm) in the electrolyte, the Ru‐catalyzed battery chemistry changes from LiThe authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


The microbial reduction of nitrate, via nitrite into gaseous di-nitrogen (denitrification) plays a major role in nitrogen removal from aquatic ecosystems. Natural abundance stable isotope measurements can reveal insights into the dynamics of production and consumption of nitrite during denitrification. In this study, batch experiments with environmental bacterial communities were used to investigate variations of concentrations and isotope compositions of both nitrite and nitrate under anoxic conditions. To this end, denitrification experiments were carried out with nitrite or nitrate as sole electron acceptors at two substrate levels respectively. For experiments with nitrate as substrate, where the intermediate compound nitrite is both substrate and product of denitrification, calculations of the extent of isotope fractionation were conducted using a non-steady state model capable of tracing chemical and isotope kinetics during denitrification. This study showed that nitrogen isotope fractionation was lower during the use of nitrite as substrate (ε = −4.2 and −4.5‰ for both treatments) as compared to experiments where nitrite was produced as an intermediate during nitrate reduction (ε = −10 and −15‰ for both treatments). This discrepancy might be due to isotopic fractionation within the membrane of denitrifiers. Moreover, our results confirmed previously observed rapid biotic oxygen isotope exchange between nitrite and water. Even though essential for plants, dissolved nitrate can be responsible for contamination of drinking water supplies from surface and groundwaters, posing a potential threat to human health10. Fortunately, it has been shown that denitrification plays a significant role in the mitigation of nitrate pollution13. Denitrification is the microbial dissimilatory reduction of NO3− to N2 via several steps, governed by different enzymes16. During this beneficial removal of NO3−, the intermediates nitrite (NO2−), a toxic component at low concentrations17, and nitrous oxide (N2O), which participates in ozone layer destruction18, are formed. Whereas the production of N2O has received considerable attention due to its role as a potent greenhouse gas, the intermediate compound nitrite (NO2−) has received much less attention since it is often assumed to be thermodynamically unstable, short-lived, and therefore not accumulating in the environment. However, nitrite has been detected at concentrations exceeding the European Water Framework Directive (EU WFD) of 0.009 mg N L−1 in urbanized rivers23. Therefore, a better understanding of its dynamics in aquatic systems is highly desirable.The growing world population requires an increased food production, necessitating the intensive use of nitrogen-containing synthetic fertilizers, which has led to nitrate pollution in many surface water and groundwater bodies2O and N2 28, anaerobic ammonium oxidation (Anammox)29, nitrifier denitrification30 and nitrification32.Nitrate reduction initially results in nitrite production. During phase 1, the “net nitrite production” phase, the NO3− (δ15N and δ18O) have been used since several decades to better distinguish the origin of NO3− 36. In addition, it provides information regarding the processes involved in its transformation, e.g. denitrification and nitrification in lakes, groundwater, riparian zones, rivers, soils, and marine environments43. Denitrification is accompanied by significant N and O isotope fractionation affecting the remaining NO3− due to reaction kinetics differences between the lighter and the heavier isotopes of N and O (14N vs. 15N and 16O vs. 18O). Enrichment in heavy isotopes of N and O in the remaining nitrate is caused by the fact that 14N and 16O react faster than 15N and 18O during denitrification44.Stable isotope techniques can provide unique insights into sources and processes regulating the N cycle in ecosystems. The nitrogen and oxygen isotope ratios of NO2) via different intermediates . The N and O isotope enrichment factors ε15N and ε18O during denitrification are usually determined by measuring variations in concentrations and isotope ratios of the remaining nitrate, while the isotopic compositions of the final product N2 or of the intermediate compounds such as NO2− and N2O are rarely analyzed. In past studies, the concentrations and isotopic compositions of nitrate and nitrite have not always been separately determined due to analytical limitations for the latter. Therefore, the “net ε15N” and “net ε18O” determined for the entire denitrification process assessed solely based on measurements of the progressively disappearing nitrate may not always be accurate.Denitrification is a unidirectional reaction in which nitrate is reduced to gaseous nitrogen (N2−) rather than nitrate as source of nitrogen for the determination of ε15N for nitrite reduction. Two of these studies were carried out in soil experiments46 another with a microbial culture43. Due to new chemical reduction analytical methods, it is now possible to determine the isotopic composition of NO2− alone, separating the isotopic analyses of NO3− from those of coexisting NO2− 50. Recent studies of the isotopic composition of nitrite focused mainly on nitrification56. A few studies mention ε15N values for nitrite reduction during denitrification. These studies were conducted using pure cultures of denitrifiers43, marine incubations42, in-situ marine environments55 or more recently abiotic denitrification57. In order to improve our knowledge on the isotope effects affecting nitrite, we investigated isotope fractionation during denitrification in laboratory experiments with environmental bacterial communities. Nitrite was either applied as a substrate or analyzed as intermediate during nitrate reduction.Only a few studies used nitrite  associated to denitrification. To our knowledge, this study is the first focusing on freshwater benthic denitrifying communities with two distinct sources of dissolved inorganic nitrogen, i.e, nitrite or nitrate.Since during denitrification and nitrate consumption nitrite is both substrate and product, it is not possible to use the Rayleigh equation to calculate concentrations and isotope compositions of the intermediate compound nitrite. For this reason, we developed the −1 g−1 were observed at high initial nitrite concentrations whereas the low initial nitrite concentration resulted in lower nitrite reduction rates of 35 nmol min−1 g−1 under the experimental conditions.In the low nitrite experiment (101 µM), concentrations decreased with time reaching 9.7 µM after 240 min Fig. . In the 15N-NO2− of the remaining NO2− increasing from −18.1‰ to −8.4‰ and from −19.6‰ to −17.6‰ for low and high treatments respectively .Similar to the experiments carried out with nitrite only, the kinetics of denitrification were investigated with solutions with low (97 µM) and high (508 µM) initial nitrate concentrations Fig. . Nitrate2− concentrations were less than 1 µM. In the low nitrate treatment, NO2− concentrations increased to 23 µM at 90 minutes and then decreased to 0.2 µM at 105 minutes , initial NO3−, the decrease in NO3− concentrations was accompanied by a corresponding enrichment of 15N in the remaining NO3−. The δ15N-NO3− of the remaining nitrate increased from 5.5 to 33.7‰ within 90 minutes and from 7.5 to 47.4‰ within 300 minutes in low and high nitrate treatments  and 7.2 × 10−5 min−1 for NO2− reduction  that are very close to the values obtained by applying the Rayleigh model to the data  experiment  than the value determined in batch experiments where only NO2− reduction occurred  that were very close to the values measured in the batch experiments where only NO3− was reduced  than that determined in batch experiments where only NO2− reduction took place (ε = −4.7‰) to obtain a good model fit for the data of the multi-step NO3− → NO2− → N2 experiment  for nitrite reduction to N2 of −4.5 ± 0.3‰ was determined. This ε15N value for nitrite reduction is on the low end of previously reported values. Chien et al.45 reported ε15N values between −33.2 and −2.9‰ in soil experiments. Martin and Casciotti (2016) calculated ε15N between −22 to −8‰ for reduction of nitrite associated with different marine strains.Even though initial nitrite concentrations affected the nitrite reduction rates, a similar isotope enrichment factor  of −21.1 and −24.4‰ are in good agreement with those found in the literature under similar conditions where denitrification occurs58. During nitrate reduction, the observed 15N enrichment of the remaining nitrate is caused by the preferential reduction of 14N-NO3− and a consequent enrichment of 14N of the product and intermediate compound NO2− during phase 1 of “net NO2− production” as shown in Fig. 2− concentrations increase and reach a maximum value during this phase 1 of denitrification  measured for nitrite reduction alone (ε = −4.2‰) in the experiment with the low nitrite treatment, and combined with the ε15N of nitrate reduction (ε = −21‰), the model did not accurately predict the measured δ15N-NO2− and constantly underestimated the measured δ15N-NO2− value  of NO2− reduction was changed to −10‰ in the model, an almost perfect fit between simulated and observed δ15N-NO2− was achieved  is higher than the ε15N obtained from the experiments with nitrite as sole source of nitrogen (ε15N = −4.2 and −4.7‰ for low and high nitrite treatment respectively).Based on the N isotope enrichment factor  than suggested by the Isonitrite model, independent of the nitrite reduction rates or nitrite concentrations. The apparent nitrogen isotope fractionation of nitrite during nitrate reduction seemed to increase with increasing availability of nitrite. The amount of nitrite available for nitrite reductase during denitrification is variable and lower than in the experiments with nitrite addition. Hence, either the variable or the lower nitrite concentrations may explain the difference in nitrogen isotope enrichment factors compared with those calculated in the nitrite reduction experiment.It is also interesting to note that the calculated rates in the model revealed that the nitrite production rate varies during nitrate reduction, as shown especially for the high nitrate treatment. The nitrogen isotope enrichment factor determined for experiments with nitrite as the sole electron acceptor was lower for both treatments , but in most cases nitrate is reduced by the membrane bound nitrate reductase Nar. The membrane bound Nar reduces nitrate to nitrite in the cytoplasm, nitrite is then transported over the cell membrane via a transporter (nark or nasH) back to the periplasm where it is reduced further by either nirS or nirK16. The affinity with the transporter might thus generate nitrogen isotope fractionation. Combined transport and limiting nitrite concentrations could thus be responsible for higher nitrogen isotope fractionation during nitrate reduction as opposed to experiments were nitrite was the sole electron acceptor.Nitrite reduction is mediated by two different types of periplasmic enzymes, the copper containing nitrite reductase (nirK or Cu-NIR) or the cytochrome cd1-containing nitrite reductase (nirS or Fe-NIR)3− concentrations is usually accompanied by a significant enrichment of 18O in the remaining nitrate44. The oxygen isotope fractionation factors of −19.6 and −21.9‰ calculated in the current study are on the high end of what has been previously reported in the literature44.During denitrification, a gradual decrease in NO15N −21.1 to −24.4‰ and ε18O −19.6 to −21.9‰) for low and high nitrate treatments with a ratio of N and O isotope enrichment factors (18ε/15ε) observed in this study of 0.9 for low and high nitrate treatments. This is in agreement with results from a number of recent isotopic denitrification studies where the observed 18ε/15ε ratio was close to 1.062. The role of the different nitrate reductases from various bacterial strains has recently been discussed potentially explaining the large range of ε15N values observed during denitrification in marine sediments42. Moreover, Frey et al.63 showed with pure cultures that the extent of N isotope fractionation during denitrification may depend on the different nitrate reduction enzymes (periplasmic Nap versus membrane bound Nar).In our study, the nitrogen and oxygen isotope fractionation during denitrification were in the same range , it was compared with the chemical kinetic constant of NO3− reduction kNO3 (Eq. + (µM−1 min−1) were converted into the units of kNO3 (min−1) by considering relation 18 and multiplying k+ (µM−1 min−1) by the concentration of H216O (55 × 106 µM). A k+ constant of 1.65 min−1 was obtained which is 5 orders of magnitude higher than the kNO3 (4.2 × 10−5 min−1). It appears thus that oxygen isotope exchange in our experiment is very rapid compared to the microbial mediated reduction of NO3−. Previous studies hypothesized that isotopic exchange can occur between the oxygen of water and the oxygen of nitrite over long time periods53. Our study confirms rapid oxygen isotope exchange between the intermediate nitrite and water. Due to the fast change of the δ18O-NO2−, it appears that isotope variations are driven not by physical parameters but are caused by bacteria activity.In the absence of published values of the kinetic constant of oxygen isotope exchange k18O-NO2− changed immediately and remained stable during nitrate reduction. Model calculations also confirmed rapid oxygen isotope equilibrium exchange between water and nitrite with a kinetic isotope constant for O isotope exchange of >5 * 10−2 kg mol−1 s−1.This study reports the N and O isotope enrichment factors associated with the multi-step biogeochemical process of denitrification. Isotope measurements and calculation of isotope fractionation based on the Isonitrite model provide insights into the apparent isotopic fractionation of nitrite during denitrification. Due to rapid isotopic exchange between the oxygen of nitrite and water, δBased on N and O isotope enrichment factors obtained from batch experiments, model simulations revealed an apparent nitrogen isotope enrichment factor for nitrite reduction (of −10‰ and −15‰) higher than those calculated for the experiments (−4.1‰ and −4.7‰) with nitrite as initial source of nitrogen with two different concentrations. The reduction of nitrate seems to drive the change of the nitrogen isotope fractionation factor affecting nitrite reduction. We hypothesize that this difference is due to variations of the amount of available nitrite. This suggests that the amount of available nitrite and transport determine the extent of nitrogen isotope fractionation during the process of denitrification.2− as the sole electron acceptor in a laboratory batch incubation experiment. Subsequently, the same experimental approach was used with NO3− as electron acceptor in order to determinate the N (ε15N) and O (ε18O) isotope enrichment during denitrification of NO3−. For both experiments, two treatments at the low (101 µM for nitrite and 97 µM for nitrate) and high end (614 µM for nitrite and 508 µM for nitrate) of concentrations measured in surface waters were used. Subsequently, a non-steady state model was developed to simulate the changes in concentrations and N and O isotope ratios of nitrate and nitrite tracing kinetics and isotope variations during reduction of nitrate concentrations associated with initial net production of nitrite and subsequent net reduction of nitrite.In order to determine the variations in nitrogen and oxygen isotope ratios, N and O isotopic enrichment factors for nitrite reduction were determined under denitrifying conditions (in the absence of oxygen) with NOorg = 8.6%; Norg = 0.8%; C/N = 14.9; δ13Corg = −24.6‰; δ15Ntotal = 2.9‰). Thirty grams of sediment were weighed in 300 ml glass bottles and 280 mL of nitrite or nitrate in milliQ water was added to create a sediment slurry. The slurries were flushed with He for twenty minutes to achieve anoxic conditions prior to crimp capping the bottles with a septum and a ring. The solutions in the first set of experiments contained 101 and 614 µM of nitrite derived from KNO2 (low and high nitrite treatments), and the second set of experiments contained 97 and 508 µM of nitrate derived from KNO3 (low and high nitrate treatments). Prior to addition, the salts were dissolved in milliQ water with a δ18O-H2O value of −10‰. Bottles were incubated under continuous agitation in order to maintain sediments in suspension and optimize denitrification  in the dark at 20 °C for 270 and 330 minutes for low and high nitrite treatments and for 180 and 420 minutes for low and high nitrate treatments, respectively. For each treatment, 14 similar slurries were prepared and considered homogeneous at the beginning of the experiment. Every 15 to 30 minutes for each treatment, a bottle was sacrificed and the slurry was centrifuged at 4000 rpm for 15 minutes. The supernatant was then filtered first through a 0.45μm glass microfiber filters (GFF-Whatman) and subsequently through a 0.2 µm nylon membrane (Whatman) to remove all sediment. Solutions were immediately preserved with mercuric chloride (0.3 ml at 5%) to avoid any residual transformation of the dissolved compounds in the solution. An aliquot of this solution was used to determine nitrite and nitrate concentrations. Ammonium concentrations were determined at the beginning and at the end of the experiments to check the ammonium budget. As ammonium consumption due to nitrification can be excluded under anoxic conditions, ammonium production should be only caused by ammonification (degradation of organic matter) or DNRA. Ammonium concentrations were below the limit of detection for all analyzed samples. Another aliquot of the solution was frozen and stored for subsequent N and O isotope analysis on nitrite and nitrate.Muddy organic sediments were collected from the Morbras Stream, La Queue en Brie, a tributary of the Marne River in the Seine watershed upstream of Paris (France). Sediments from this stream had been used in a preliminary test and showed significant nitrate reduction potential (data not shown). The land use in the upstream catchment is mainly crops. The sediment was homogenized in the laboratory prior to the experiments. Contents and isotopic compositions ot total organic carbon and total nitrogen were determined by EA-IRMS  to obtain nitrite or nitrate and ammonium concentrations respectively with a precision of the measurement ±0.3 µM.64. Nitrate was reduced to nitrite in an activated column of cadmium. Yields of nitrate reduction were calculated using concentration measurements with a yield of reduction >95%.Nitrogen and oxygen isotope ratios of nitrate and nitrite were determined separately following a modified protocol of McIlvin and Altabet2− at the column outlet, which is composed of reduced NO3− and the initial NO2− . Subsequently, the NO2− (i.e. NOx− and NO2−) was further reduced to nitrous oxide by sodium azide solution. To achieve this, each sample was diluted with a salt solution (NaCl = 35.5 g l−1) to obtain 1 µM of NO2− concentration in 15 ml and transferred into a 20 ml glass vial. The vials were crimp-sealed with Teflon-backed silicone septa  and aluminum caps. The sodium azide solution, which consists of a 40 ml mixture of 20 ml of azide  and 20 ml of acetic acid , was flushed with He for 30 min to avoid any N2O pollution from the buffer. In order to reduce the NO2− to N2O, 0.8 ml of azide solution was injected in each vial, and vials were placed in a water bath at 30 °C for one hour. The yield of conversion was better than 95%.When both nitrite and nitrate concentrations were not negligible, the liquid samples were diluted with a salted buffer to obtain a concentration of 20 µM of NO2O samples were measured with an isotope ratio mass spectrometer  in continuous-flow mode with a purge-and-trap system coupled with a Finnigan GasBench II system . Results are reported in the internationally accepted delta notation in ‰ with respect to the standards air for δ15N and Vienna Standard Mean Ocean Water (V-SMOW) for δ18O, respectively. Nitrate and nitrite reference materials subject to the same analytical procedures were used to calibrate the isotopic composition of N2O . The precision was ±0.3‰ for δ15N-NO2− and ±0.5‰ for δ18O-NO2−.The isotope compositions of all Nx− (NO2− + NO3−) and NO2− alone, the δ15N-NO3− and δ18O-NO3− were calculated as follows:3−] and [NO2−] are the concentrations of NO3− and NO2−, respectively.Finally, using the isotopic compositions of NO15N-NO3− was ±0.5‰ and ±0.8‰ for δ18O-NO3−.Based on this calculation, the precision for δ15N and δ18O-NO2− (t0) are the N and O isotope ratios of the substrate at time zero, and C0 and Ct are the concentrations of nitrite at time zero and t, respectively. This relationship was used to determine the N and O isotope enrichment factor of nitrite reduction by plotting measured values of δ15N(t) and δ18O(t) against ln Ct/C0. The same approach was used to determine the N and O isotope enrichment factor for denitrification for the experiments where nitrate was the sole source of nitrogen.The nitrite experiment in a closed system without any renewal of substrate ensured Rayleigh conditions and therefore the N and O isotope enrichment factors (ε) were calculated using the following equations:2− when these two processes occur concurrently in a batch reactor. Although our batch experiments alone are sufficient to derive enrichment factors (ε) for the separately tested NO3− = >NO2− and NO2− = >N2 steps, a model was developed to calculate the isotope enrichment factors when these processes occur simultaneously in a batch reactor. The model simulates the multi-step denitrification process and considers the presence of the intermediate product (NO2−) as well as the N and O isotopic compositions of nitrite and nitrate. The model is constrained with experimentally determined concentrations and N and O isotopic ratios of nitrite and nitrate. The model was used to estimate instantaneous isotopic enrichment factors associated with NO3− reduction to NO2− and NO2− reduction to N2. Moreover, the model was used to investigate the kinetics of oxygen isotope exchange between water and NO2−.As nitrite is both a substrate and a product during denitrification of nitrate, Rayleigh conditions and equations are not applicable to determine the N and O isotope enrichment factor associated with the reduction of NO3− and NO2− concentrations are expressed in µM, Rx (x = NO3− or NO2−) is the nitrite production or reduction (reaction) rate in µM min−1, kx is the kinetic constant in min−1 and nx is the reaction order.The model considers that denitrification is a two-stage process:18O3: 18O16O16O and 18O2: 18O16O.To simulate both N and O isotope ratios simultaneously, three isotopic species must be considered for each chemical species, with the total chemical concentration being equal to the sum of the isotopic species:15N and the δ18O values of NO2− and NO3− can be calculated as follows (equations shown for NO3):18O/16O ratio in the SMOW standard and NSTD is the 15N/14N ratio of air.From these isotopic species, the δ3− and NO2− reduction for the different isotope species have to sum up to the rate of the total chemical species calculated with Eqs. 3− reduction:The rate of reaction for NO15εNO3 and 18εNO3, respectively):2− and H2O takes place according to the following isotopic exchange reaction:16O2 = 14N16O2 + 15N16O2, the rates are in µM min−1, and the forwards (k+) and backwards (k−) kinetic constants (in µM−1 min−1) are related to the equilibrium constant of the isotope exchange reaction a follows:i refers to the initial, equilibrium concentration of isotopic species.The expressions for the rates of the isotopic species are obtained by solving a system of three equations comprising equation 14 and equations 12 and 13 written in terms of rates rather than concentrations, and introducing isotopic enrichment factors for N and O isotopes  the oxygen and nitrogen isotope compositions of NO3− and NO2−, and (4) the δ18O of water. Unknown (fitting) parameters are: (1) the kinetic constants for NO2− production, for NO2− consumption and for isotope exchange between NO2− and H2O, and (2) the oxygen and nitrogen isotope fractionation factors in NO2− production and NO2− consumption.To summarise, known model input parameters are: (1) the simulation time (that corresponds to the duration of the experiment), (2) the initial concentrations of NO2− and NO3−.The model was implemented with the software Mathematica version 9 (Wolfram Research). Model outputs are changes of nitrite and nitrate concentrations and nitrogen and oxygen isotope compositions for both NO


In the remaining five cases sudden increases and decreases were made in the operating conditions of the PV module. Finally, the theoretical data of the PV system are presented, which can be used as a reference to analyze the data obtained with the two controllers in the six cases. Data are provided in This article presents the data of a PV system using fuzzy logic and the perturb and observe (P&O) algorithm, under sudden changes in the solar irradiance and the operating temperature of a PV module. The mathematical modeling of the PV module, the dc-dc converter and the fuzzy and P&O controllers were discussed in our previous work entitled "Fuzzy logic based MPPT controller for a PV System"   Specifications tableValue of the data•With these data, readers can make comparisons with other algorithms used to track the maximum power point in a PV system. This saves time for comparisons in different weather conditions •These data are useful for the implementation and experimental evaluation of MPPT controllers with high efficiency •The data obtained can be used to train MPPT controllers using different intelligent control techniques such as neural networks.•These data are useful to analyze in detail the behavior of a PV system under sudden changes in the solar irradiance and the operating temperature of the PV module.1This paper presents the numerical data obtained from the modeling and simulation of a PV system consisting of a 65 W PV module, a buck converter, a 12 V battery, a standard P&O controller and a fuzzy controller; under sudden changes in the solar irradiance and the operating temperature of a PV module. The simulations were performed on a computer with Windows 10 pro operating system, Intel Core i5-6500 processor and 8 GB RAM. 22 and constant operating temperature of 25 °C. Case 1: Controllers were evaluated for standard test conditions, with a constant solar irradiance of 1000 W/m2 and 1000 W/m2 with increments of 200 W/m2. Case 2: Controllers were evaluated for a constant operating temperature of 25 °C and a solar irradiance signal that increases in different instants of time. This signal is between 200 W/m2 and 200 W/m2 with decreases of 200 W/m2. Case 3: In this case the controllers were evaluated for a constant operating temperature of 25 °C and a solar irradiance signal between 1000 W/m2 and an operating temperature that increases in different instants of time with increments of 25 °C. This signal is between 0 °C and 100 °C. Case 4: In this case the controllers were evaluated for a constant solar irradiance of 1000 W/m2 and an operating temperature that is decremented from 100 °C to 0 °C with decreases of 25 °C. Case 5: Controllers were evaluated for a constant solar irradiance of 1000 W/mCase 6: Controllers were evaluated for changes in operating temperature and solar irradiance. For the two signals, dynamic profiles with sudden increases and decreases at different instants of times were used. In order to obtain the voltage, current and output power of the PV system for sudden changes in operating conditions, dynamic profiles were used for the solar irradiance and the operating temperature of the PV module. Using a fuzzy controller and the P&O algorithm, six (6) test cases were defined:Finally,


From SdH oscillations, the evidence of Dirac fermions with π Berry phase is found and the experimentally determined two main Fermi wave vectors are correlated to two surface Dirac cones (buried one inside the other) of layered topological insulator PbBi4Te7. We have also found evidence of conductance fluctuations, the root mean square amplitude of which is much higher than the usual universal conductance fluctuations observed in nanometer size sample. Calculated autocorrelation functions indicate periodic unique fluctuations may be associated with the topological surface states in the compound.We present a detail study of Shubinikov-de-Haas (SdH) oscillations accompanied by conductance fluctuations in a mesoscopic topological insulator PbBi Angle resolved photoemission spectroscopy (ARPES) measurement in prototype Bi2Se33 and several other TI single crystal samples6 demonstrated that an odd number of massless, spin helical Dirac cones (DC) are present at their surface states. Transport8 and scanning tunneling microscopic9 measurements corroborate the spins at the surface states are locked perpendicular to the momenta. In various studies, investigations related to the topological character of these fascinating materials through transport measurements provide some important and novel information e.g., weak antilocalizatoin11, Shubnikov-de-Hass (SdH) oscillations14, quantum fluctuations16, etc. Other than the above phenomena, Checkelsky et al.17 have observed a unique magneto fingerprint signal in micro-meter size Ca doped Bi2Se3 single crystal flake at low temperature with angle dependent magnetoresistance (MR) measurement and argued that the novel fluctuations are related to the spin degree of topologically protected surface states. One of the crucial limitation in transport studies to investigate the surface states is the interference from metallic bulk states which hinder the perfect spin polarization at the surface states and the large number of the carrier density that shifts the Fermi level towards the bulk conduction band instead pinning to Dirac point (DP). Moreover, topological insulator materials exposed to the environment also manifest a shift in Fermi level towards the bulk conduction band19. The surface exposure can be reduced by using inert gas atmosphere, photo resist layer etc., but these involves complications. Therefore, to overcome surface exposure as well as to get maximum surface polarization it is instructive to choose the material very carefully. In search of new TI materials, numerous binary as well as ternary chalcogenides have been investigated along with different doping elements and by applying gate voltages to tune the band gap and the Fermi level26. Surface and bulk sensitive ARPES studies show that Bi2Se3 has the largest band gap27 whereas, PbBi2Te4 shows the maximum spin polarization (70%)6. ARPES studies of PbBi4Te7 confirm the existence of two Dirac cone, DC1 and DC2, one buried inside the other28. The unit cell of PbBi2Te4 is made by insertion of PbTe layer in Bi2Te3, consist seven atomic monolayers and similarly, hexagonal unit cell of PbBi4Te7 single crystal is made by five atomic layers of Bi2Te3 block and seven-atomic layers of PbBi2Te4 block, separated by van-der-Waals gap. As crystals used to cleave at the van-der-Waals gap, so, the crystal is terminated either at 5 layers of Bi2Te3 block or at 7 layers of PbBi2Te4 block. In the case of 5 layers of Bi2Te3 terminated surface, the surface sates of PbBi4Te7 distributed not only at the top 5 layers, but also extended upto next 7 under layers of PbBi2Te4; unlike 7 layers terminated surface where the surface Dirac cone located solely at the 7 layer28. Therefore deep buried topological surface states one inside the other in PbBi4Te7 receive physical protection without the need of any external efforts. Moreover, due to the existence of large spin polarized surface states in Pb and Bi based alloys we may able to observe clear signature of the magneto finger print as observed by Checkelsky et al.17, motivated us to study topological insulator PbBi4Te7. Reduction of the flake thickness of TI single crystals is a powerful tool to elucidate the surface states lead to the fabrication of nano device of PbBi4Te7 single crystal with thickness 60 nm and length of few micrometer.The study of the interplay between symmetry and phase transition in condensed matter physics has achieved great importance for several decades. Recently the existence of a new class of material called topological insulators (TIs) has been predicted, where symmetry alone cannot describe its chiral spin polarization at the edge states4Te7 single crystal, synthesized by modified Bridgman method. We present the observations of weak antilocalizatoin, Shubnikov-de-Hass (SdH) oscillations and unique conductance fluctuations. Experimentally determined π Berry phase confirmed the Dirac nature of the carriers. The estimated Fermi wave vectors correspond to the two surface Dirac cones (buried one inside other) of PbBi4Te7. We don’t observe the signature of universal conductance fluctuations (UCF) as expected for the nano device, instead unique conductance fluctuations are observed.Herein, we present for the first time an exploration of the magneto-transport properties of nano device made from mechanically exfoliated flake of PbBi4Te7 single crystal. To investigate the growth direction of the crystal, flakes were exfoliated along the lengths of the ingots, for each flake we observed sharp c-axis reflections from their Bragg’s planes, inferred the high crystalline quality of the single crystal. To investigate the phase purity of the single crystal, the exfoliated flakes were pulverized in to polycrystalline powder. Rietveld refinement29 of the corresponding powder XRD is illustrated in the inset (left panel) of Fig. 4Te7 belong to P-3m1(164) space group with a = b = 4.455 Å and c = 24.155 Å respectively, good agreement with literature32. The absence of parasitic peaks confirms the monophasic nature of the single crystal. Refined lattice parameters are modeled33 and are illustrated in the inset (middle panel) of Fig. 4Te7 consists of 12 atomic layers of Pb, Bi and Te. In the unit cell, Te-Bi-Te-Bi-Te, five layers atomic block (5 LB) of Bi2Te3 is separated from seven layer block (7 LB) of Te-Bi-Te-Pb-Te-Bi-Te i.e. PbBi2Te4 via van-der-Waals gap, the alternate stacking of the 5 LB and 7 LB along c direction make the 3D structure of PbBi4Te7. At room temperature two distinct Raman modes at 104 cm−1 and 124 cm−1 are identified for the PbBi4Te7 single crystal for parallel to crystallographic c axis measurement. The mode positions are verified by Lorentzian line shape fitting of the modes as illustrated in the inset (right panel) of Fig. Figure 4Te7 nano device having flake thickness ~60 nm [see Fig. c axis and perpendicular to the direction of the sensing current. Figure xx) as a function of temperature for the nano device at various applied DC magnetic fields  parallel to crystallographic c axis. As the temperature decreases, resistance starts to decrease up to ~50 K, with saturating nature bellow it, the behavior is typical for a metallic sample. Temperature dependency of Rxx can be best described by: 0 is the residual resistance of the crystal. The exponential and quadratic terms represent electron-phonon and electron-electron interactions, respectively. R0 = 0.01042 Ω, β = 0.00916 Ω, θ = 527.56 K and γ = 7.779 × 10−9 Ω K−2 corresponded to the best fit for B = 0 T curve. The fitted parameters indicate the conduction at higher temperatures is assisted by electron-phonon scattering mainly whereas, the electron-electron interaction can be neglected. The determined phonon frequency is −1. In the presence of magnetic field (B = 2 T and 5 T), the R0 value does not change considerably; R0 = 0.0104(7) Ω for B = 2 T and R0 = 0.0104(4) Ω for B = 5 T, respectively as illustrated in Fig. 35.Figure see Fig. , illustrxx(B)) at different temperatures with a magnetic field applied perpendicular to the basal plane of the crystal. The common features of the curves are fluctuations along with oscillations in the entire range of magnetic field and non-saturating enhancement of magnetoresistance upto highest applied field (8 T). The reproducibility of the oscillatory features of the curves at different temperatures make it distinct from the random noise. To investigate the effect of magnetic field on the resistance of PbBi4Te7 nano device, magnetoresistance curves are smoothed using first order polynomial36. The subtraction of the smoothed background of the actual signal gives the fluctuations, the nature of which discussed in detail later. The smoothed curves as illustrated in Fig. xx. For better contrast, derivative of the longitudinal conductivity fSdH be the frequency of oscillations, B be the applied magnetic field and γ be the phase factor. The fast Fourier transform (FFT) of the oscillations at 2 K reveles most prominent frequencies at 276 and 120 T possessing highest amplitudes  correspond to the observed frequencies at 276 T and 120 T (the two most prominent frequencies) are determined from Onsagar’s relation: −1  and 0.060 Å−1 , respectively. The estimated KF value for 158 T (at 10 K), ~0.069 Å−1, is nearly equal to the KF2 observed at 2 K. In order to determine the effective mass of electrons, amplitude variation of oscillations with temperature is fitted with Lifshitz-kosevich (LK) factor37: ωc is the cyclotron frequency and related to the effective mass (m*) as: e (KF2) and 0.079 me (KF1) at field values ~7 T and ~5.7 T, respectively, me is the free electron mass. The corresponding Fermi velocity, vF5 ms−1 (KF2) and 8.80 ×105 ms−1 (KF1), respectively. The estimated position of Fermi energy level from Dirac point of KF2 is 4Te7 28. The ratio of the wave vectors  corresponding to the observed frequencies at 276 T and 120 T is close to the reported ratio  of DC2 and DC1 at binding energy 300 meV28, indicating KF1 and KF2 correspond to topological surface states of DC1 and DC2, respectively. ARPES reveled shape of the constant energy contours for DC1 shows more energy dependency than DC2 i.e. DC1 shows more warping effect of Fermi surface over DC228. The warping of Fermi surface leads to peak splitting as illustrated in Fig. τ) is determined from the field dependence of SdH oscillations for a fixed temperature i.e. from Dingle analysis of two different regions (corresponds to each beat) of 2 K data. Figure 34 D = ABsinhD); A is the amplitude of oscillation determined by averaging the vertical distance of crests from the adjacent valleys and α = −1. The Dingle temperatures (TD) correspond to DC1, DC2 are 22 and 27 K, respectively, obtained from the slopes of the ln(D) vs τ) from TD values −14 s (DC1) and 5.98 × 10−14 s (DC2), respectively. The obtained mean free paths lSdH(=vFτ) of the carriers are 48.4 nm (DC1) and 22 nm (DC2), respectively. The mobility μSdH2 V−1 s−1, respectively. The high value of μSdH and low effective masses of DC1, DC2 confirms their surface origin. The estimated μSdH value for DC2 is small compared to DC1 and quite above the bulk mobility of prototypes TIs38, reflects the DC2 is deep buried inside the DC1. At 2 K the estimated surface carrier densities ns12 cm−2 for DC1 and 6.58 × 1012 cm−2 for DC2.Figure des Fig.  and at 1des Fig. . Dividin39. To explore the possibility whether the fluctuations observed are UCF, we plot absolute δG δG is background subtracted conductance. Zoomed in regions of Fig. δG at different temperatures. It can be seen that the absolute conductance fluctuations are nominally depends on magnetic field in the entire range of magnetic field at different temperatures. To quantify the observed fluctuations, δGrms [see Fig. 40δGrms is observed, indicating the surface origin of conductance fluctuations41 The decreasing fluctuations amplitude upon increasing the temperature contemplates us to consider the transport model of UCF proposed by Lee et al.42. But, the very high value of δGrms compared to the universal value 17 C = s) in FFT of the correlation function obtained are 1.5 T and 3 T (2nd harmonic of 1.5 T), respectively at 2 K. For 10 K it, is at 1.8 T [see inset of Fig. ϕ value at each temperature are smaller than the flake thickness (60 nm) is another evidence of the surface origin of conductance fluctuations. The oscillations decay with increasing field, which is same as observed by Checkelsky et al.17, and exactly opposite to the observation we have for the SdH oscillations. The characteristics length, lϕlϕ ~ T−0.25 [see Fig. 2Se3 by Checkelsky et al.17, arising due to spin degree of topological surface states, though the angle-dependent conductivity measurement in future is necessary to relate the observed fluctuations with spin degree of surface states. Here it is noteworthy that the large fluctuations superimposed on the background of magnetoresitivity data are not affecting the topological surface states with non-trivial π Berry phase, attributed the robustness of its topological phase.We now focus our attention on the conductance fluctuations obtained by subtraction of the smoothed background of the actual magnetoresistance signal. At low temperatures, universal conductance fluctuations (UCF) are particularly significant in single crystalline nanoflakes when the thickness of the flake is comparable to the phase coherence length of the charge carriers. The UCF arises from interference of electrons paths around the defect centers upon sweeping the magnetic fieldsee Fig.  is detersee Fig. . In the 4Te7 device. The Shubinikov-de-Haas oscillations corroborate two unambiguous transport signature, first, it provides evidence of surface Dirac fermions with π Berry phase and secondly, the determined two wave vectors by FFT analysis are closely related to the two surface Dirac cones (buried one inside the other) of the layered topological insulator PbBi4Te7. Though, it is natural to observe Universal conductance fluctuations in nanometer size sample yet, the root mean square amplitude is inconsistent with UCF. The field variation of autocorrelation function of the fluctuations indicate the observed fluctuations may like the magnetofingerprint observed in Ca doped Bi2Se3 crystal17.In conclusion, we report the Shubinikov-de-Haas oscillations and conductance fluctuations in a mesoscopic topological insulator PbBi4Te7 was grown by melting the stoichiometric homogeneous mixture of Bi, Pb and Te of analytical grade with purity >99.99% in an evacuated quartz tube at 1223 K, the ampule was then cooled down to 893 K at a rate of 2 Khr−1. followed by furnace cooling to room temperature. The crystal quality and the phase identification were carried out through room temperature x-ray diffraction (XRD) data collected by a 9 kW Rigaku X-ray diffractometer, operated at 3 kW, equipped with Cu-Kα radiation in the 2θ range of 10° to 90°. Chemical homogeneity was confirmed by the elemental analysis of the energy dispersive spectrum (EDS) of X-rays recorded with Carl Zeiss EVO10 SEM system with an Oxford X-ray detector. Room temperature Raman spectra were collected with Technos STR-500 micro-Raman spectrometer equipped with 532 nm diode laser source having spectral resolution of 1 cm−1. To avoid surface oxidation, the laser power was kept at <2 mW. Nano device was made using an exfoliated flake of uniform thickness ~60 nm (confirmed by atomic force microscopy) on clean Si/SiO2 wafers (with SiO2 layer thickness ~300 nm) by e-beam lithography. Electrical contact pads were made by thermal deposition of Cr (4 nm)/Au (60 nm) bi-layers followed by lift-off where exposures of the flakes to air were minimized by using inert gas atmosphere. In plane electronic transport measurements were carried out with AC transport option of Physical Property Measurement System (PPMS)  using a dc magnetic field applied perpendicular to the sensing current and along the crystallographic c (00l) direction.Single crystal of layered chalcogenide PbBi