DA) and the resultant formation of powerful oxidants such as the hydroxyl radical (•OH) through ‘Fenton chemistry’ in the presence of iron within dopaminergic neurons are thought to contribute to the damage of cells or even lead to neuronal degenerative diseases such as Parkinson's disease. An understanding of DA oxidation as well as the transformation of the intermediates that are formed in the presence of iron under physiological conditions is critical to understanding the mechanism of DA and iron induced oxidative stress. In this study, the generation of H2O2 through the autoxidation and iron‐catalyzed oxidation of DA, the formation of the dominant complex via the direct reaction with Fe(II) and Fe(III) in both oxygen saturated and deoxygenated conditions and the oxidation of Fe(II) in the presence of DA at physiological pH 7.4 were investigated. The oxidation of DA resulted in the generation of significant amounts of H2O2 with this process accelerated significantly in the presence of Fe(II) and Fe(III). At high DA:Fe(II) ratios, the results from this study suggest that DA plays a protective role by complexing Fe(II) and preventing it from reacting with the generated H2O2. However, the accumulation of H2O2 may result in cellular damage as high intracellular H2O2 concentrations will result in the oxidation of remaining Fe(II) mainly through the peroxidation pathway. At low DA:Fe(II) ratios however, it is likely that DA will act as a pro‐oxidant by generating H2O2 which, in the presence of Fe(II), will result in the production of strongly oxidizing •OH radicals.The non‐enzymatically catalyzed oxidation of dopamine  have previously been thought to be generated through the interplay between dopamine (DA) and iron, contributing to damage to cells and, potentially, leading to neuronal degenerative diseases such as Parkinson's disease. Our results suggest that DA plays a dual role as high DA/Fe(II) ratios prevent Fe(II) from reacting with the generated H2O2 thereby reducing •OH generation, whereas low DA/Fe(II) ratios enhance •OH generation as a result of reaction of unbound Fe(II) and H2O2 produced via both autoxidation and iron‐catalyzed oxidation of DA. It is reported that both the dopamine‐derived quinones and concomitant generation of ROS are associated with neurodegenerative disorders or even lead to cancer and cell death as a result of their ability to cause damage to lipids, proteins, and DNA ethane, DA) is a well‐known neurotransmitter in humans  and to form complexes with both Fe(II) and Fe(III) ‐DA complexes and the oxidation of Fe(II) through the direct DA‐iron reaction. Based on the experimental data, a detailed kinetic model capable of describing the complicated interplay between both Fe(II) and Fe(III) species, the iron oxide ferrihydrite, DA, oxygen, and the intermediate oxidation products is developed.In this study, the kinetics and mechanism of the interaction of DA with Fe(II) and Fe(III) are investigated at physiological pH of 7.4 over a range of metal to ligand ratios under both deoxygenated and oxygen‐saturated conditions, with emphasis on the generation of HAll analytical grade chemicals were purchased from Sigma‐Aldrich  (or as otherwise stated) and were used without further refinement. All solutions were prepared using 18 MΩ cm ultrapure Milli‐Q water (MQ) . All glassware was acid washed in 5% v/v HCl for at least 1 week before use. Stock solutions were kept in dark bottles and were refrigerated at 4°C when not in use. All experiments were conducted under dark conditions and performed at a controlled room temperature of 22 ± 0.6°C.3, and 10 mM 3‐(N‐morpholino) propanesulfonic acid (MOPS). Low concentrations of MOPS were found to have a negligible effect on the rates of Fe(II) oxidation 2?6H2O) in 5 mM HCl. Concentrated stock solutions of 10 mM Fe(III) [using ferric chloride hexahydrate (FeCl3·6H2O)] and 10 mM DA were prepared weekly in 10 mM HCl. The working stock solutions of Fe(II), Fe(III), and DA were diluted from the concentrated stock solutions daily in 10 mM HCl solution. The acidity of both concentrated stock and working stock solutions was sufficient to avoid significant oxidation of Fe(II) and DA and precipitation of Fe(III) on the time scale of interest and yet low enough to minimize any pH change that might occur on addition of the stock to experimental solutions. A stock solution of 20 mM H2O2 prepared by dilution of a nominal 30% w⁄w H2O2 solution was used for calibration of the H2O2 measurements. The nominal 30% w⁄w solution was standardized by UV spectrophotometry at 240 nm  and 50 KU/L horseradish peroxidase (HRP) were prepared in MQ water as described in Bader et al. (A concentrated Fe(II) stock solution (5 mM) was prepared by dissolving ferrous ammonium sulfate hexahydrate  with a maximum absorbance at 562 nm and molar absorptivity of ε562 nm = 30 000/M/cm  was quantified spectrophotometrically using the FZ method (Stookey 2O2 formed during the course of oxidation of DA [both in the absence and presence of Fe(II) or Fe(III)] was quantified using the photometric DPD method  and Fe(II) given the relative strengths of the DA and DTPA complexes with Fe(II)  were undertaken with calibration curves constructed by linear regression of the calibration data. Interference arising from the presence of low concentration of DA was found to be negligible for the measurement of H2O2 in the presence of phenanthroline (data measured is not provided here). Interference arising from the presence of Fe(III) and DA was found to be negligible under the experimental conditions investigated here .The Hmono‐complex (denoted hereafter as FeIIIDA), the bis‐complex (denoted hereafter as FeIIIDA2), and the tris‐complex (denoted hereafter as FeIIIDA3) with the dominant species being pH and concentration‐dependent ], addition of Fe(III) into solutions containing DA is expected to produce three different complexes: the IIIDA2 was determined spectrophotometrically by measuring the absorbance at 570 nm using a Cary 60 spectrophotometer with baseline correction at 850 nm  was developed under deoxygenated conditions to prevent any oxidation and transformation of the complex. Briefly, the solutions were sparged for 1 h using a special gas mixture of 297 ± 6 ppm CO2 in argon (BOC) prior to the addition of DA. The solution was then bubbled for another 10 min before the addition of Fe(III). Continuous sparging was maintained during the course of the deoxygenated experiment. The molar absorptivity of FeIIIDA2 was calculated to be 3312/M/s which is quite similar to that of the bis‐complex Fe(III)‐catechol. It should be noted that although the tris‐complex FeIIIDA3 has a higher absorptivity than that of the bis‐FeIIIDA2 (data measured under deoxygenated conditions is not given here), it would be expected to have minimal effect on the measurement of FeIIIDA2 in view of its very low concentrations under the experimental conditions investigated here. In addition, Fe(II)‐DA complexes are colorless (data measured under deoxygenated conditions is not provided here) (Tyson and Martell IIIDA2.The concentration of FeKnowledge of the speciation of both Fe(II) and Fe(III) in the presence of various concentrations of DA was utilized to determine the major iron‐DA complexes present at particular concentrations of added iron and DA. Fe(II) and Fe(III) speciation calculations were undertaken using the program Visual Minteq  or Fe(III) were modeled using a simplistic approach originally developed by Rose and Waite . In thiset al. 2O2 and the formation of FeIIIDA2 , the major Fe(III)‐DA complex under the conditions of interest. The interplay of Fe(II) and DA in the presence of O2 was finally fitted because of the complicated transformation between Fe(II) and Fe(III) with the oxidation of Fe(II) also accounted for in the model fitting. In the proposed reaction scheme  was generated in the presence of 20 μM DA.As shown in Fig. 2, after the addition of Fe(III), very rapid initial formation of FeIIIDA2 followed by a gradual increase in its concentration was observed over the first hour. The concentration of FeIIIDA2 was also found to increase with increase in the DA:Fe(III) concentration ratio. At later times, the increase in FeIIIDA2 concentration was negligible. It was also observed that not all the added Fe(III) was complexed to DA presumably because of the rapid precipitation of Fe(III) at this pH.As shown in Fig. 2, FeIIIDA2 complex also formed immediately after the addition of Fe(III) into DA containing solutions with a subsequent increase in concentration with increase in DA concentration. However, as shown in Fig. 2, the concentration of FeIIIDA2 complex continued to increase significantly over time, resulting in a substantially higher concentration of FeIIIDA2 complex in the presence of O2 at the conclusion of the 2 h study than was the case in the absence of O2.As shown in Fig. O2∙−to the observed increase in concentration of FeIIIDA2 in the presence of O2, 50 KU/L SOD was added to the solution to capture any O2∙−produced. As shown in Fig. IIIDA2.To assess the possible contribution of et al. As shown in Fig. IIIDA2 complex were rapidly formed on addition of Fe(II) to solutions containing DA but, unlike the case in which Fe(III) was added to DA‐containing solutions where there was both a rapid initial formation of FeIIIDA2 complex and an ongoing more gradual formation, the concentration of FeIIIDA2 complex gradually decreased following the initial rapid formation on Fe(II) addition. The rapid, initial formation of FeIIIDA2 complex  was converted to FeIIIDA2) confirmed that Fe(II) was first complexed by DA to form FeIIDA (reaction 22) with this Fe(II) complex then transforming to the Fe(III) complex through reaction with oxygen (reaction 23). However, the concentration of FeIIIDA2 decreases over time, particularly in the presence of low concentrations of DA  more than doubling over the course of the 2 h study. The initial rapid formation of H2O2 may be attributed to the rapid conversion of Fe(II) to Fe(III) in the presence of high DA concentration as discussed above. However, once Fe(II) was exhausted, the rate of H2O2 generation after the initial stage  from the FeIIIDA2 complex and the subsequent rapid re‐oxidation of Fe(II) to Fe(III) with concomitant reduction in oxygen to superoxide . Results of DA concentration dependency studies  via the short‐lived cyclization product, leukoaminochrome, which immediately undergoes redox exchange with the remaining DAQ yielding DAC and regenerating DA  is the dominant species and is formed through the direct reaction between iron and DA via two sequential steps: an initial rate‐limiting formation of the FeIIIDA complex (either through the direct reaction between Fe(III) and DA or the oxidation of the FeIIDA complex following its formation through the direct reaction between Fe(II) and DA) followed by replacement of one Fe(III)‐coordinated H2O by another DA molecule.At the physiological pH investigated in this study, with the co‐existence of iron and DA, the blue‐purple colored FeIIIDA2 following its initial formation. Internal electron transfer via the direct reaction between the Fe(III) metal center and the coordinated ligand DA might be expected to be particularly important and will result in loss of FeIIIDA2 with concomitant release of Fe(II) and DA•− (reaction 15) with the rate of decay depending upon the extent of coordination with the mono‐complex being the least stable and the tris‐complex being the most stable  was reported for the monocatecholate complex of Cu(II) (Kamau and Jordan bis‐complex FeIIIDA2 (k15 = 7.26 × 10−5/s) reported in this study is undoubtedly in line with this argument. However, if decay via internal electron transfer within the complex is the dominant process controlling the fate of FeIIIDA2, then this should result in a continuous decrease in concentration of FeIIIDA2 following its initial rapid formation both from Fe(III) and Fe(II). This phenomenon, however, is not apparent from the results shown in Figs IIIDA2 are involved including, potentially, the DA‐induced dissolution of any precipitated Fe(III) [initiated by the adsorption of DA to the amorphous ferric oxide (AFO) surface followed by release of FeIIIDA to solution (reaction 10)] and/or the reductive dissolution of any Fe(III) precipitate [initiated by the adsorption of DA to the AFO surface followed by electron transfer within the Fe(III)‐DA surface complex and subsequent release of Fe(II) (reaction 11)] or even the O2∙− mediated dissolution of AFO. However, as removal of O2∙− by adding SOD did not make much difference in the formation of FeIIIDA  and DA•− (reaction 18). The re‐oxidation of the released Fe(II) by DA•− exerts significant influence on the formation of FeIIIDA2, especially at the later stage of the reaction. In contrast to the deoxygenated condition, in the presence of O2, both the rapid oxidation of Fe(II) at pH 7.4 (reaction 19), which directly contributes to the generation of Fe(III) (and thus of FeIIIDA2), and the oxidation of DA•− by O2 drive the reductive dissolution reaction to the right, making the reductive dissolution of AFO be the main factor contributing to the increase in concentration of FeIIIDA2.Model simulations suggest that, under deoxygenated conditions, the reductive dissolution of AFO is not as effective as the ligand (DA)‐induced process for the overall formation of FeIIIDA2 complex increased significantly with increasing DA concentration. The increase in concentration of the FeIIIDA2 complex with increasing DA concentration in the absence and presence of O2 is reasonable since a higher concentration of DA would be more effective in preventing Fe(III) precipitation and accelerating Fe(II) oxidation. At high DA concentration , presumably, any Fe(III) existing in the solution or released from the complex (as a result of dissociation of the complex) or any Fe(II) existing in the solution or released from the complex  would be rapidly re‐bound (in the case of Fe(II), following reoxidation) by DA. In comparison, at low DA concentrations, Fe(III) hydrolysis may compete with complexation resulting in the formation of AFO rather than FeIIIDA2. Similarly, the oxidation of any existing or released Fe(II) may result in formation of AFO rather than FeIIIDA2 if there is insufficient DA to compete with the hydrolysis reaction. The gradual decrease in the concentration of FeIIIDA2 observed at 20, 10, and 1 μM DA /Fe(II) half‐reaction. Details of the expected change in the standard redox potential of the Fe(III)/Fe(II) half‐reaction are shown in the supplementary information (Text S1).In the presence of DA, the oxidation of Fe(II) becomes more thermodynamically favorable as the formation of the colorless mono‐complex with Fe(II)  with H2O2 generated from both DA autoxidation and iron‐catalyzed oxidation of DA may result in the generation of •OH through ‘Fenton chemistry’. As shown in Figs 2O2 that was produced in the process may react quickly with the remaining Fe(II) via the Fenton reaction to generate Fe(III) and the powerful oxidizing species •OH. However, in the presence of high DA concentrations, as shown in Figs IIDA and then FeIIIDA2 immediately upon addition. Thus, even though significant concentrations of H2O2 were generated in the presence of Fe(II), the Fenton reaction is unlikely to be important under these conditions as the majority of Fe(II) was locked in the FeIIIDA2 complex before it had a chance to react with the large amount of H2O2 generated during this process. This is also in agreement with the proposed protective role of phenol‐type compounds in view of their ability to bind Fe(II) with resultant prevention of the reaction of Fe(II) with H2O2, •OH production and subsequent cell and DNA damage  (with associated •OH production) would almost certainly outcompete the oxygen‐mediated oxidation of Fe(II) (a pathway which does not result in production of •OH).As a transition metal with several spin states, iron can catalyze the oxidation of DA and accelerate the generation of Het al. et al. et al. As the presence of excess neural labile iron is detrimental and may cause biological damage as a result of the generation of ROS, a chelation strategy may help to reduce the oxidative damage associated with regional iron deposition. Indeed, a variety of chelators have been investigated for the rectification of disorders related to iron‐overload over the past few decades   the effect of DA on Fe(II) oxidation  foII) Figs  and Fe(III) Figs  to solution Fig. . Rate coO2∙− and H2O2 is energetically unfavorable because E0(DA•−/DA) = 1.353 V > E0(O2∙−/H2O2) = 1.05 V and E0(H2O2/•OH) = 0.39 V  indicated that the oxidation of DA by O2∙− was only important when the rate constant of this reaction is over 1 × 103/M/s. Therefore, the oxidation of DA by both O2∙− and H2O2 were not considered in the reaction scheme. In contrast, the oxidation of DA•− by O2∙− is much more energetically favourable because E0(O2∙−/H2O2) = 1.05 V > E0(DAQ/DA•−) = 0.15 V (Pham and Waite bis‐FeIIIDA2 complex in the presence of O2 at the later stage of the reaction. The rate constant used in this study for the redox exchange between leukoaminochrome and DAQ (k5 = 5.30 × 106/M/s) was adopted from a previous dopaquinone study in view of the similarity in structure between dopamine and dopaquinone . This result suggests that the formation of melanin from DAC was not important in the early stage of the DA oxidation and thus was not included in the proposed model.From a thermodynamic perspective, the oxidation of DA by both r was quite sensitive to change in the magnitude of each rate constant. On the other hand, only an upper value of 7.26 × 10−5/s  was deduced for internal electron transfer between Fe(III) and the coordinated DA within the bis‐FeIIIDA2 complex (reaction 15).In the presence of DA, as shown in Fig. S4c, both the ligand‐induced dissolution of AFO as a result of DA adsorption onto the solid surface (reaction 10) and the DA‐induced reductive dissolution of AFO as a result of internal electron transfer (reaction 11) were important processes as the relative residual IIIDA complex from the direct reaction between Fe(III) and DA (reaction 12), the initial concentrations of the precipitate  and of FeIIIDA2 thus were used to estimate crudely the formation rate constant of the FeIIIDA complex by using the published Fe(III) precipitation rate constant (5.0 × 106/M/s)  (including both dissolved and precipitated ferric species).To have a better understanding of the formation rate constant of the rate‐limiting FeT), i.e. [DA] ≈ [DA]T, substituting [Fe(III)I] = [Fe(III)]T − [FeIIIDA] and [Fe(III)] = [Fe(III)]T − [FeIIIDA] − [AFO], the rate laws may be written as:k12′=k12[DA]T. Solution for this linear OED is of the form ]T = 5 μM, [DA]T = 50 μM, k9 = 5.0 × 106/M/s and [FeIIIDA]eq ≈ [FeIIIDA2] = 1.23 μM to equation 5 gives k12 = 1.43 × 105/M/s. Similarly, at 100 μM DA and 5 μM Fe(III), substituting [Fe(III)]T = 5 μM, [DA]T = 100 μM, k9 = 5.0 × 106/M/s and [FeIIIDA]eq ≈ [FeIIIDA2] = 1.70 μM into equation 5 gives k12 = 1.07 × 105/M/s. These approximate values were used as a constraint for the model fitting of k12.At 50 μM DA and 5 μM Fe(III), substituting [Fe(III)]5/M/s was deduced for formation of the mono‐complex FeIIIDA (reaction 12). This value is consistent with the formation rate constants proposed for the complexation of a range of natural organic ligands with Fe(III) (Rose and Waite bis‐complex FeIIIDA2 from FeIIIDA and DA (reaction 13) is assumed to be similar to the rate constant for water‐loss from Fe(OH)(H2O)52+ of 4.50 × 105/M/s  is almost three orders of magnitude less than that with dissolved inorganic Fe(III) (of 1.50 × 108/M/s). Given that the reactivity of O2∙− with AFO is strongly dependent upon the age and structure of AFO  and DAQ) with a rate constant varying from 10−2 to 106/M/s  did not vary the relative residual r with this result suggesting that this reaction was unimportant and, as such, was not included in the model.Based on the analysis above and sensitivity analysis , a best‐fit rate constant of 2.50 × 102/M/s  is consistent with values proposed for Fe(II) complexation with a range of naturally occurring organic compounds  and formation rate constants  is expected to form complexes with DA more slowly than with Fe(III)  and Fe(III). However, because of the considerable insensitivity of the relative residuals to variation in the rate constants of the reactions between organically complexed Fe(II) with DA•−, O2∙− and H2O2  and the reduction in organically complexed Fe(III) with O2∙− , these rate constants were assumed to be of similar order of magnitude to that for reaction of these species with inorganic iron. In contrast to the reactions with O2∙− as shown in Fig. S4h, reduction in both inorganically and organically complexed Fe(III) with DA•− is not important given the insensitivity of the relative residuals to variation in the rate constants (from 1 to 109/M/s) of these reactions. Therefore, the reduction in inorganically and organically complexed Fe(III) by DA•− was not considered in the proposed reaction scheme.It is clear from the results obtained that oxidative intermediates DA2, a considerable amount of H2O2 can be generated during the autoxidation of DA with this process significantly accelerated in the presence of both Fe(III) and Fe(II). In the presence of Fe(III), rapid formation of the bis‐complex FeIIIDA2 at physiological pH followed by slow internal electron transfer within the complex results in production of Fe(II) and DA•− which, in the presence of oxygen, are oxidized with concomitant formation of superoxide. The superoxide so formed then either disproportionates or is further reduced, resulting in production of H2O2 in addition to that produced by DA alone. Observation of the initial increase in concentration of FeIIIDA2 in the absence and presence of oxygen suggested that mobilization of iron from amorphous ferric oxide by DA through both ligand‐induced dissolution and reductive processes was operating at the physiological pH used in the studies described here. As DA is more effective in inducing AFO dissolution in the presence of oxygen, we conclude that the reductive mechanism of DA‐mediated AFO dissolution predominates. Observation that addition of SOD exerted a minor influence on the formation of FeIIIDA2 suggested that O2∙− ‐induced Fe(III) dissolution from AFO may not be as important as previously thought.The results of this study show that, in the presence of OIIDA followed by reaction with O2 resulted in conversion of almost all Fe(II) present in this complex into the more stable FeIIIDA2 complex at high DA concentrations. While this direct reaction between Fe(II) and DA resulted in rapid generation of a substantial amount of H2O2, the deleterious effect of reaction of this H2O2 with Fe(II) (the ‘Fenton reaction’) was avoided as the majority of the Fe(II) present was rapidly converted to FeIIIDA2 via an oxygenation pathway. In comparison, at low DA concentrations, a slow oxygenation rate of Fe(II) coupled with a rapid depletion of generated H2O2 in the presence of excess iron suggested that Fe(II) might react with H2O2 to form strongly oxidizing •OH. Even though the oxidation of Fe(II) by O2 may still be important, the presence of additional H2O2 as a result of the presence of DA provides a pathway for the formation of strongly oxidizing •OH.In the presence of Fe(II), instantaneous formation of the mono‐complex Fe2O2 through the oxidation of DA in the absence and presence of Fe(II) and Fe(III), the formation of FeIIIDA2 complex in the presence of Fe(III) and Fe(II) in both oxygenated and deoxygenated environments and the oxidation of Fe(II) in the presence of DA over a range of DA/iron ratios at physiological pH. While there are many additional factors that would need to be considered in extending this model to description of in vivo intracellular processes, this model should assist in better understanding the anti‐oxidant and pro‐oxidant properties of intracellular DA, especially with regard to its role in the transformation of intracellular iron. In addition, the mechanistically based kinetic model developed here should aid in predicting the appropriate dosages of iron chelators that could be used in clinical trials aimed at examining the efficacy of chelation therapy to treatment of neurodegenerative disorders such as Parkinson's disease.A relatively simple kinetic model has been developed and found to satisfactorily describe the overall kinetics of generation of HTable S1. Stability constants for Fe(II) and Fe(III) speciations at 25°C and I = 0.Figure S1. Measured absorbance of H2O2 at 551 nm with baseline corrected at 690 nm in 0.1 M NaCl at pH 7.4 in the absence of Fe(III) and DA  and in the presence of 5 μM Fe(III)  and 20 μM DA . Error bars are standard errors from triplicate measurements.Figure S2. Speciation of Fe(III)  and Fe(II)  over the pH range 4.0–8.5 in solutions containing varying concentrations of DA with [Fe(II)]0 = 5 μM and [Fe(III)]0 = 5 μM. LogC denotes the log concentration of individual ferrous and ferric species.Figure S3. Measured absorbance of FeIIIDA2 complexes (a) and calibration curve for qualification of FeIIIDA2 complexes measured at 570 nm with baseline corrected at 850 nm (b) in 0.1 M NaCl at pH 7.4 in the presence of 400 μM DA under deoxygenated condition.Figure S4. Sensitivity analysis for different fitting reaction rate constants /Fe(II) half‐reaction.Click here for additional data file.


The appearance of these excitons is attributed to the change in many‐body interactions that occurs alongside the interfacial orbital hybridization and spin–orbit coupling brought about by the excitonic effect propagated from the substrate. This has further led to the formation of a Fermi‐surface feature at the interface. The results provide an atomic‐scale understanding of the heterointerface between monolayer‐TMDs and perovskite oxide and highlight the importance of spin–orbit–charge–lattice coupling on the intrinsic properties of atomic‐layer heterostructures, which open up a way to manipulate the excitonic effects in monolayer TMDs via an interfacial system.The exciton, a quasi‐particle that creates a bound state of an electron and a hole, is typically found in semiconductors. It has attracted major attention in the context of both fundamental science and practical applications. Transition metal dichalcogenides (TMDs) are a new class of 2D materials that include direct band‐gap semiconductors with strong spin–orbit coupling and many‐body interactions. Manipulating new excitons in semiconducting TMDs could generate a novel means of application in nanodevices. Here, the observation of high‐energy excitonic peaks in the monolayer‐MoS The Raman spectrum of MoS2/Al2O3 in Figure E2g1 ‐mode (≈386 cm−1) due to opposing vibrations of in‐plane Mo and S‐atoms, and A1g‐mode (≈406.1 cm−1) due to opposing vibrations of two out‐of‐plane S‐atoms. The wavenumber difference between these two modes is ≈20.1 cm−1, consistent with previously reported for monolayer‐MoS2/Al2O3.2 is of high quality. Upon transferring monolayer‐MoS2 onto STO substrate, E2g1 ‐mode is blueshifted while A1g‐mode remains unchanged—indicating an increase in compressive strain between monolayer‐MoS2 and STO,2 as demonstrated later. The compressive strain is also demonstrated in PL spectra in Figure 2/STO are significantly lower than the respective spectra of MoS2/Al2O3. This contrast in intensity highlights the strong influence optical interference and absorption effects have on the Raman and PL intensities of monolayer‐MoS2 on different substrates. Further observation of interfacial effects such as charge transfer at the MoS2/STO interface will be presented and discussed later.Large‐area monolayer‐MoS1(ω)+iε2(ω)) of MoS2/STO and MoS2/Al2O3 (inset: absorption coefficient) as measured by spectroscopic ellipsometry. This experimental technique utilizes the analysis of multilayer models2 . Significant optical renormalization is observed in Figure 2 and absorption spectrum of MoS2/STO while that of MoS2/Al2O3 falls to zero. This indicates the presence of carriers2/STO but absent in MoS2/Al2O3. The estimated effective number of carriers in monolayer‐MoS2 and interface is calculated from the optical conductivity spectra at energy range up to 1.5 eV,2 unit formula . These carriers can be attributed to complex lattice–charge–orbit coupling as discussed later. At intermediate‐energy range, three main peaks labeled A (≈1.88 eV), B (≈2.04 eV), and C (≈3 eV) are observed for MoS2/Al2O3. They are related to the excitonic transitions2 is transferred onto STO, consistent with the ≈40 meV blueshift of excitonic peak in PL spectra in Figure 2/STO is always higher than that of MoS2/Al2O3. This spectral weight is transferred from the higher energy range (>3.3 eV). According to Zaanen–Sawatzky–Allen theory,2/STO, significantly different from MoS2/Al2O3.Figure 2/STO. This peak is absent from monolayer‐MoS2 on Al2O3 and other substrates.2/STO interfacial layer by means of interfacial orbital hybridization that couples with the strong electron–hole interaction .Interestingly, at high‐energy range, a new peak D appears at ≈3.8 eV for MoSn Figure f. To fur2.22 and STO . Due to the formation of Mo—O bonds (to be demonstrated later), the oxygen atoms from atmosphere are introduced at the sulfur‐vacancy sites during their growth and transfer processesFigure2.2 with and without the STO substrate are shown in Figure 2/STO in Figure First‐principles calculations are performed to further understand the interfacial interaction between monolayer‐MoS2/STO . As temperature decreases, peak D sharpens and is redshifted. Recent studies have shown that excitonic effects in STO due to strong electron–hole interactions can be observed in its optical spectra.Δα = α−α) where there is similar trend of sharpening and redshifting with decreasing temperature. This feature has been attributed to Wannier‐like excitons in STO.2/STO and the corresponding spectra of STO at the same energy as well as the GW‐BSE calculation provide clear evidence that the excitonic wavefunctions of STO substrate extend into monolayer‐MoS2 through interfacial orbital hybridization.Figure 2. This implies an interfacial Ti—S orbital hybridization. A previous study also shows weak interfacial adsorption with interfacial orbital hybridization.Figurez‐orbitals of S‐atoms interact directly with Ti‐3d orbitals, leading to a significant orbital resonance between −4 and −0.5 eV. More specifically, the density of states calculation of the Ti and S atoms as displayed in Figure z orbitals and the t2g orbitals belonging to Ti‐3d near the valence and conduction band edges. This larger overlap implies a stronger interaction between these two sets of aforementioned orbitals within this energy region. The energy level of the t2g‐orbitals belonging to Ti‐3d is lower than the eg‐orbitals and they are located closer to the Fermi level. This suggests a stronger interaction between the S‐pz orbital and the Ti‐3d t2g orbitals. This will be experimentally demonstrated in our X‐ray linear dichroism (XLD) absorption measurement as discussed later  compares the absorption spectra of the previous MoS2/STO at 300 K and the ex situ spectrum in its current state. The absorption spectrum of the aged MoS2/STO sample shows an intensity reduction slightly below that of the original 300 K measurement. This result further indicates the continual O substitution that reduces the Ti‐S hybridization at the MoS2/STO interface. While there is an ageing process taking place in the MoS2/STO with O substitution into the S‐vacancies, the sample remains rather stable even under atmospheric condition.Results from first‐principles calculations further confirm the presence of interfacial orbital hybridization. As shown in Figure Interestingly, the exciton peak splits into two peaks, D and D′, at low temperature  and ≈3.80 eV region. Besides, these spectral differences are consistently present with repeated measurements at different temperatures . The spectral differences between monolayer‐MoS2/STO and STO are studied by a fitting analysis using two Gaussian lineshapes as displayed in Figure 2 (≈1 nm) being much thinner than the STO substrate, the PL signal suggests an influential effect of the high‐energy interfacial excitonic responses on the radiative properties of monolayer‐MoS2.To further evaluate the excitonic effects brought about by the interfacial dynamics of monolayer‐MoS2.4Figure2/STO, MoS2 on Cu, and bulk MoS2 using synchrotron‐based photoemission spectroscopy (PES). According to previous photoemission spectroscopy studies,z2 and S3pz (and O2pz) orbital; peak b is ascribed to O2p and Ti3d bands intermixed with MoS2 p‐d valence band; peak c is ascribed to the p‐d covalent bonding band between O and Mo‐atoms intermixed with MoS2 p‐d valence band.2/STO interface shown in the close‐up of valence band spectrum. This is consistent with the optical result in Figure 2/STO. It implies that the excitonic effect propagation from the STO substrate also has an appreciable impact on the weak electronic correlations in the MoS2 monolayer. The increased weak electronic correlations in MoS2 may affect the Fermi surface instability that is induced by symmetry breaking of the MoS2 film that manifests itself in the form of the Fermi‐surface feature. This feature is not observed for MoS2 films on other substrates. The MoS2/STO valence band maximum  is shifted to a higher BE as compared to MoS2/Cu but is lower than that of bulk‐MoS2. Figure 2/STO compared to bulk‐MoS2 and MoS2/Cu. The shifts in work function and VBM* in MoS2/STO are possible effects of interfacial orbital hybridization, charge transfer, and lattice distortion. As noted in Figure 2. Since the peak of bulk MoS2 is broad due to point defects, surface roughness, impurities, and vacancies, to ensure consistency, the work function is derived by intersecting the extrapolated leading edge with the background baseline (details are provided in the Supporting Information). Schematics of the band offsets are depicted in Figure 2/STO can be harnessed to increase the efficiency of nanoelectronic devices.Interestingly, an anomalous Fermi‐surface feature  measurements were performed. It is noted that monolayer‐MoS2 is too thin to yield decent Mo and S XAS signals for effective characterization. Therefore, we turn instead to the characterization of the Ti L and the O K‐edges in the XAS measurements. This indirect but effective characterization methodology that bypasses the restriction of monolayer‐MoS2 being too thin serves to provide essential information on the formation of the Mo—O bonds. The Ti L and the O K‐edges of the bulk STO are further used as a reference to confirm the formation of Mo—O bonds that will be discussed subsequently. Polarization‐dependent Ti L‐edges XAS for MoS2/STO and bulk‐STO are shown in Figure4+ and Ti3+ in Figure j→Ti3d transitions. The spectra show four features attributed to spin–orbit‐coupling splitting (2p3/2 and 2p1/2) and crystal‐field splitting (eg and t2g). With the multiplet calculations as references, the increased intensities at ≈458 and ≈464.5 eV  model is sensitive to few monolayers. This indicates the presence of Ti3+ ions at the interface and the onset of charge transfer. This is consistent with the estimated effective electron number (≈0.01 per MoS2 unit formula) from ellipsometry measurement. This electron‐transfer process is further supported by polarization‐dependent Ti L‐edge XLD shown at the bottom of Figure 2g band . The spectra difference (I(E||c)‐I(E||a), linear dichroism) implies a preferential occupancy of the 3dxy orbital. It demonstrates the presence of mixed Ti3+/4+. The positive integrated XLD implies a larger out‐of‐plane absorption—more out‐of‐plane (dxz/yz) empty states while the transferred electrons are in the in‐plane orbitals (dxy) .Polarization‐dependent O K‐edge XAS of the MoS2/STO, it needs to combine both Ti L3,2 and O K‐edges XAS spectra. In Figure 2/STO, particularly at 530–532 eV (O2p‐Ti3d hybridization states), is significant. While the change of XAS yield at Ti L3,2 edges between bulk STO and MoS2/STO in Figure To further analyze the interfacial charge transfer and the electronic structure in MoS2g‐O2p (≈530.4 eV) and Ti3deg‐O2p (≈532.8 eV) peaks are nearly polarization independent, the new O2p‐Mo4d peak (≈531.9 eV) is polarization dependent. The spectra for the in‐plane (90°) and out‐of‐plane differ due to the crystallinity of Mo—O.2 reported in a previous study.2 are partially occupied by the O‐atoms , which are shown in Figures S7 and S8 (Supporting Information). This is to provide further verification that interfacial hybridization and the formation of the new high‐energy exciton still take place even with the stacking of MoS2 and STO in different configurations as long as there are strong interfacial interactions between the MoS2 and STO substrate. It also shows that the onset of the new exciton is attributed to the unique interfacial hybridization and it is weakly correlated to the stacking configuration between the TMDs and the STO substrate.To further substantiate the agreement between experimental and theoretical studies, we perform further first‐principles studies involving monolayer‐MoS32/STO heterointerface, it provides a deeper insight into the mechanisms leading to the new excitonic and electronic effects. The study of the interfacial phenomena is based on a comprehensive methodology that comprises temperature‐dependent spectroscopic ellipsometry, photoluminescence, X‐ray absorption and photoemission spectroscopies, and first‐principles calculations. Detailed study of these anomalous excitons uncovers new insights into existing theories, which is related to the interband transitions, and underlines many of the intriguing physical properties, thereby exposing them to further examination.By unraveling the complex spin–orbit–charge–lattice coupling at MoS4Sample Preparation: The monolayer‐MoS2 was synthesized on a sapphire (Al2O3) surface by the CVD method using MoO3 and S powders as the reactants.2 was then transferred to TiO2‐terminated (001)‐STO substrate and Cu foil using polymethylmethacrylate. The TiO2‐terminated STO substrate was obtained through ultrasonic treating in buffered hydrofluoric acid solution and annealing in a furnace at 950 °C for 1.5 h in air with a ramping up rate of 5 °C min−1 and ramping down rate of 3 °C min−1. In this study, the TiO2‐terminated STO substrate was chosen because the 3d electrons of the transition metal Ti‐atoms play an important role in generating new and exotic properties in both the STO bulk materials and in the formation of other heterostructure systems.2 was of high quality before and after the chemical transfer process. These experiments confirmed that the samples retained their high quality even after the transfer. Besides, the method of chemical transfer of monolayer‐MoS2 was preferred over the direct synthesis of MoS2 on SrTiO3 using the CVD process as the latter technique would introduce oxygen vacancies in the SrTiO3 substrate at high temperature in vacuum. It was demonstrated that sulfur vacancies were inevitably present in CVD‐fabricated monolayer‐MoS2 and that they were further generated during the transfer processes.2/STO heterostructure. This was a clear demonstration of O substitution into S‐vacancies in our monolayer‐MoS2 samples.Spectroscopic Ellipsometry Measurements and Absorption Coefficient: A J. A. Woollam Co., Inc spectroscopic ellipsometer with photon energy of 0.6–4.5 eV was used to measure the ellipsometry parameters Ψ (the ratio between the amplitude of p and s‐polarized reflected light) and Δ (the phase difference between p and s‐polarized reflected light) in a high vacuum chamber with a base pressure of 1 × 10−9 mbar. The substrate layers (bulk SrTiO3 or Al2O3) were also measured under the same conditions. The absorption coefficient, α, of MoS2 monolayer was extracted from parameters Ψ and Δ utilizing an air/MoS2/STO (or Al2O3) multilayer model, where monolayer‐MoS2 consisted of a homogeneously uniform mediumEffective Electron Number: Taking into consideration our optical data, by analyzing the optical conductivity spectrum of MoS2/STO at room temperature,m0 is taken as the free‐electron mass and V is the unit cell volume V = 0.054 nm3 of monolayer‐MoS2. The Neff(ω) is proportional to the number of electrons involved in the optical excitations up to ω. The optical conductivity at energy between 0 eV (estimated as 0) and 0.6 eV was estimated using a linear interpolation. The estimated effective number of carriers Neff (1.5 eV), calculated from the optical conductivity spectra at energy range up to 1.5 eV,X‐Ray Absorption Spectroscopy (XAS): The O K‐edge absorption spectra in the energy range 520–580 eV and Ti L‐edge absorption spectra in the energy range 450–490 eV were obtained using linearly polarized X‐ray absorption spectroscopy from the Surface, Interface and Nanostructure Science (SINS) beamline at Singapore Synchrotron Light Source (SSLS), using a TEY detection method. The incidence angle (90‐θ) of X‐rays refers to the normal of the sample surface, which was varied by rotating the polar angle of the sample. The spectra were normalized to the integrated intensity between 565 and 580 eV for O1s spectra and between 475 and 490 eV for Ti2p spectra after subtracting an energy‐independent background.X‐Ray Linear Dichroism (XLD) Measurements: For Ti, the 3dxy and 3dxz/yz orbitals had lobes pointing parallel and perpendicular to the ab‐plane, respectively. XLD could probe the occupancy of the 3dxy and 3dxz/yz orbitals using linearly polarized light aligned to the out‐of‐plane  and in‐plane  directions : The room temperature photoemission data were taken in an ultrahigh vacuum chamber with a base pressure of 1 × 10−10 mbar at the SINS beamline of SSLS. The photon energy of 60 eV was used to probe the valence band spectra. The work function was measured using 60 eV photon energy with a −7 V applied bias. The spectra were collected at normal emission using a VG Scienta R4000 analyzer and normalized by photon current. The binding energy is referred to the Fermi level of a sputter‐cleaned gold foil in electrical contact with the sample.Atomic Multiplet Calculations: The simulations were performed using CTM4XAS software with octahedral (Oh) crystal field 10 Dq = 1.85 eV, Hubbard value Udd = 1.7 eV, and charge‐transfer energy Δ = 1.2 eV. The Slater integrals were taken to be 80% of the Hartree–Fock values, which represent the atomic values.4+‐multiplet states, calculations showed an overall spectral shift of the Ti3+‐multiplet states to lower photon energies , (√3 × 2) MoS2 supercell was placed on the (√2 × √2) STO (001) substrate with six atomic layers, in which tensile strain (about 14.2% tensile strain on the b‐direction) was applied on the STO as the electronic properties of monolayer‐MoS2 are sensitive to the external strain. The differential charge density was visualized by an isosurface value of 1.0 × 10−3 e Å−3. For the perfect monolayer‐MoS2 on STO (001) substrate, the interaction between them was found to be weak as the calculated adsorption energy is 37 meV Å−2 and layer spacing is about 2.62 Å, similar to that of monolayer‐MoS2 on the HfO2 or graphene on the STO substrate.2 during the transfer process. Thus, O defects were considered in MoS2 monolayer by substituting an S atom in the interface structure, which was related to surface oxygen density of 1/8. The Mo—O bond length is shorter than that of Mo—S bond. Hence, compared to the perfect monolayer‐MoS2, local compressive strain was expected in monolayer‐MoS2 with such O‐defects—consistent with Raman and PL data  and shorter interfacial spacing (2.58 Å). It was noted that in experiment the oxygen concentration in MoS2 might not be so high. Due to highly demanding GW‐BSE calculations, a lager supercell could not be used, but it is thought that the physics discussed here would not be affected significantly by this relatively small interface supercell. The adsorption of monolayer MoS2 on the STO was weak as the calculated adsorption energy was about 44 meV Å−1. The weak interfacial interaction is also confirmed by the unnoticeable interfacial charge transfer in Figure on STO 00,  with the Perdew–Burke–Ernzerhof format exchange‐correlation functional and the projector augmented wave potentials.k‐point mesh was used for sampling the first Brillouin zone.2 monolayer. Van der Waals correction to the dispersion energy was included by using Grimme's DFT‐D3 method.U = 4.0 eV was applied to d‐orbital of Ti atoms in order to take into account the orbital‐dependent Coulomb and exchange interaction. Besides, dipole correction was included in the calculations to minimize the errors introduced by the periodic boundary conditions. In all calculations, the energy and force on each atom were converged to be smaller than 1.0 × 10−6 eV and 0.01 eV Å−1, respectively. The bottom two layers of STO were fixed during the structural relaxation.2, a more advanced GW‐BSE method was used to calculate the optical spectra of monolayer‐MoS2/STO interface and monolayer‐MoS2 structures, where the many‐body effects were well described. In GW‐BSE calculations, more than 300 empty bands and a plane‐wave cutoff of 200 eV were included for the summation of the imaginary dielectric function. A total of 30 highest valence bands and 30 lowest conduction bands were included as a basis for the excitonic eigenstates. With the presence of the STO substrate, the electron–hole interaction was partially screened, leading to a binding energy estimated at ≈0.52 eV by comparing the energy position difference between the first transition peaks from the GW‐RPA and GW‐BSE calculations in Figure 2 monolayer were not fully compensated by the O atoms, making the MoS2 monolayer electron‐doped. While in the simulations, due to highly demanding GW‐BSE calculation, only a relatively small size of MoS2/STO supercell was considered, in which one S vacancy was compensated by an O atom, the whole system was still semiconducting. Split peaks A and B in Figure Due to the large excitonic effects in monolayer‐MoSThe authors declare no conflict of interest.SupplementaryClick here for additional data file.


The incorporation of boron dopant in TiO2 ETL not only reduces the hysteresis behavior but also improves PCE of the perovskite device. The simultaneous improvements are mainly ascribed to the following two reasons. First, the substitution of under‐coordinated titanium atom by boron species effectively passivates oxygen vacancy defects in the TiO2 ETL, leading to increased electron mobility and conductivity, thereby greatly facilitating electron transport. Second, the boron dopant upshifts the conduction band edge of TiO2, resulting in more efficient electron extraction with suppressed charge recombination. Consequently, a methylammonium lead iodide (MAPbI3) photovoltaic device based on B‐TiO2 ETL achieves a higher efficiency of 20.51% than the 19.06% of the pure TiO2 ETL based device, and the hysteresis is reduced from 0.13% to 0.01% with the B‐TiO2 based device showing negligible hysteresis behavior.Perovskite solar cells (PSCs) have witnessed astonishing improvement in power conversion efficiency (PCE), more recently, with advances in long‐term stability and scalable fabrication. However, the presence of an anomalous hysteresis behavior in the current density–voltage characteristic of these devices remains a key obstacle on the road to commercialization. Herein, sol–gel‐processed mesoporous boron‐doped TiO 2 (B‐TiO2)Mesoporous boron‐doped TiO is demonstrated as an improved electron transport layer (ETL) for perovskite solar cells for the reduction of hysteresis. The incorporation of boron dopant in TiO2 ETL not only reduces the hysteresis but also improves device performance. Consequently, a methylammonium lead iodide photovoltaic device based on B‐TiO2 ETL achieves a promising efficiency of 20.51% with negligible hysteresis. This encourages more effort to explore efficient ETLs with higher conductivity and less defects.Since the first demonstration of perovskite solar cells (PSCs) by Miyasaka and co‐workers in 2009, tremendous attention has been devoted to this field due to their rapidly increased power conversion efficiency (PCE) and potentially low fabrication cost.2 ETLs, surface modifications and doping strategies have been developed.2 ETL to mitigate interfacial recombination and improve the interface binding in planar PSCs. They got hysteresis‐free planar PSCs with enhanced efficiency and stability.2 nanoparticles, and achieved high efficiency with suppressed hysteresis by applying NH2‐TiO2 as the ETL.2. To tackle this problem, a variety of dopants, such as Al, Y, Nb, and Li, have been investigated to manipulate carrier behavior.Voc) of PSCs is also reduced due to the downward shift of conduction band minimum (CBM) of TiO2. It should be noted that downshift of TiO2 CBM does not always cause Voc drop as it is not the only factor affecting the voltage. Voc can be enhanced or maintained when these dopants passivate electronic trap sites in the bulk and at the surface of the TiO2. By comparison, dopants such as Nb that upshift the CBM of TiO2 should be more desirable. Nevertheless, the hysteresis effect has not been resolved well for these dopants . Doping in the TiO2 lattice can produce considerable lattice disorder and defects, thus causing the decrease of electron mobility and the increase of parasitic recombination.2 ETL simultaneously to achieve efficient electron extraction and hence balanced carrier transport. Previously, we have shown that boron (B) doping could enhance the electron injection from excited dye molecules to TiO2 photoanode, resulting in the increased PCE of dye‐sensitized solar cells. Recently, Tian et al. reported that a low concentration of boron doping could enhance the carrier mobility and electrical conductivity of TiO2.2 (abbreviated as B‐TiO2) is expected to be a desirable ETL for PSCs to synchronously improve the PCE and significantly reduce the hysteresis.In order to improve the charge transport properties of TiO2 as an improved ETL for PSCs. B‐TiO2 ETL exhibits upshifted CBM, increased electron mobility and conductivity, and reduced electron traps. Our density functional theory (DFT) results show that B doping leads to stronger binding at the ETL/perovskite interface with a total binding energy of −5.41 eV for TiO2/perovskite interface and −6.63 eV for the B‐TiO2/perovskite interface. These calculation results are further supported by steady‐state photoluminescence (PL), time‐resolved photoluminescence (TRPL), and thermal admittance spectroscopy (TAS) results. The incorporation of B dopant in TiO2 ETL not only reduces the hysteresis behavior but also improves Voc, short‐circuit current density (Jsc), and fill factor (FF) of perovskite device. As a result, PSC based on B‐TiO2 ETL demonstrates a higher efficiency of 20.51% than 19.06% of an undoped TiO2 ETL based device, and the hysteresis is reduced from 0.13% to 0.01% with the B‐TiO2 based device showing negligible hysteresis behavior.Herein, we demonstrate sol–gel‐processed mesoporous B‐TiO22.12 and B‐TiO2 nanoparticles were synthesized by a facile sol–gel method and subsequent hydrothermal treatment, as schematically illustrated in Figure S1a in the Supporting Information.2 and B‐TiO2  samples are well‐dispersed nanoparticles, and both have a narrow distribution in the average size of ≈21 nm . The B‐TiO2 sample  shows an unchanged interplanar spacing of 0.358 nm for the (101) plane of anatase TiO2. Scanning electron microscopy (SEM) images of both TiO2 and B‐TiO2  films exhibit typical porous surface structure, which will facilitate the infiltration of perovskite precursor solution. Furthermore, atomic force microscopy (AFM) and corresponding cross‐section SEM views show that B doping leads to a slight decrease of the root‐mean‐square (RMS) roughness compared to pristine TiO2 .Pure TiO2 and B‐TiO2 samples. The peak positions  of B‐TiO2 did not show any discernible difference to pure anatase, and the values of the full width at half maximum (FWHM) of these two samples were almost the same (0.522 for TiO2 and 0.524 for B‐TiO2). To confirm this, XRD patterns were measured of TiO2 and B‐TiO2 powder scraped from corresponding sintered films and XRD patterns of TiO2 and B‐TiO2 films deposited on fluorine‐doped tin oxide (FTO)/cp‐TiO2 are also presented . It can be inferred that the doping level of 2% had insignificant effect on the crystal structure of TiO2.X‐ray diffraction (XRD) patterns  confirm that a pure anatase phase (JCPDS No. 21‐1272) formed after calcination and there was no observable structural difference between TiO2 lattice, X‐ray photoelectron spectroscopy (XPS) measurements were conducted. Figure2 sample, with a single peak centered at about 192.1 eV, which is consistent with literature values.2 counterpart does not have a corresponding peak. Figure 1/2 and Ti 2p3/2 peaks of the TiO2 sample at binding energies of ≈464.3 and ≈458.6 eV, respectively, which are typical for the Ti—O bonds in TiO2.2, the Ti 2p3/2 peak is shifted slightly higher to 458.7 eV, resulting from charge transfer through B doping, suggesting that the B dopant was incorporated in the TiO2 lattice without the formation of a separate phase. In Figure 2 sample further suggests the linkage of the Ti—O—B bond in B‐TiO2. Moreover, based on the XPS fitting results, the amount of B in the B‐TiO2 sample is evaluated by an elemental sensitive factor method,To verify the form of incorporation of B in the TiO2.2Figure2 and B‐TiO2 films. Both of them display high average transmittance in the visible region, demonstrating excellent optical quality. The UV–vis absorption spectra in Figure S5 in the Supporting Information show a slight red shift with B doping. Figure Ahν)2 versus hν, which were obtained from the absorption spectra of TiO2 and B‐TiO2 samples. The bandgap (denoted by Eg) of the TiO2 is 3.81 eV while that of B‐TiO2 is 3.78 eV, which is consistent with that reported in the literature.High optical transmittance of the charge transport layer is crucial for maximizing the photocurrent of perovskite devices. 2 films before and after doping. Figure S6a,b in the Supporting Information shows that the secondary electron cutoffs of both TiO2 and B‐TiO2 is 16.14 eV corresponding to the work function of ≈5.06 eV. Figure S6c in the Supporting Information shows valence band spectra, from which the energy differences between the valance band (denoted by Ev) edges and Fermi level (denoted by Ef) of TiO2 and B‐TiO2 were found to be 2.84 and 2.78 eV, respectively. The Ev of TiO2 and B‐TiO2 was evaluated to be −7.90 and −7.84 eV, respectively. Finally, combining with optical bandgaps, the conduction band (denoted by Ec) edges of TiO2 and B‐TiO2 were determined to be −4.09 and −4.06 eV, respectively. In Figure S6d in the Supporting Information, we illustrate that the Ec of B‐TiO2 is shifted by around 0.03 eV toward the vacuum level, which is beneficial for enhancing Voc.[qv: 14b]Ultraviolet photoelectron spectroscopy (UPS) measurements were also performed to ascertain if there is any change of the energy levels of the TiO2 and B‐TiO2 samples were measured by the space charge limited current (SCLC) model employing an electron‐only device configuration . The detailed calculation process is provided in Note S1 in the Supporting Information. Figure J–V) curves of the electron‐only device fitting the Mott–Gurney law.2 is 3.30 × 10−5 cm2 V−1 s−1, while it increases to 1.69 × 10−4 cm2 V−1 s−1 after B doping, about five times higher than the former. The higher electron mobility would contribute significantly to the electron transport, and therefore higher Jsc and FF in the perovskite cells. To determine direct current conductivity (σ0) of TiO2 and B‐TiO2 samples, we fabricated FTO/cp‐TiO2/TiO2 or B‐TiO2/Ag devices, and the current–voltage (I–V) characteristics are shown in Figure I–V plot, I = σ0AD−1 V, where A is the sample area (9 mm2) and D is the film thickness (180 nm). The calculated values of the conductivity are 6.00 × 10−3 and 1.81 × 10−2 mS cm−1 for TiO2 and B‐TiO2, respectively, which is consistent with the variation trend of the electron mobility .In addition to energy band alignment, the carrier mobility also affects the charge transfer process. The electron mobilities of TiO2 films, the dark I–V curves of TiO2 and B‐TiO2 devices  were measured, as shown in Figure VTFL) was determined by the trap state densitye is the elementary charge of the electron (1.6 × 10−19 C), L is the TiO2 film thickness (180 nm), ε0 is the vacuum permittivity (8.854 × 10−12 F m−1), εr is the relative dielectric constant of TiO2 (55),nt is the trap state density. With the averaged VTFL value , the trap state density nt in B‐TiO2 was calculated to be 6.39 × 1016 cm−3. For comparison, the calculated trap state density in the TiO2 was 1.27 × 1017 cm−3, which is almost two times higher than in the B‐TiO2. Clearly, B doping effectively passivates electron traps, resulting in the increased electron mobility.Low mobility of the ETLs indicates more trap states, which would lead to charge accumulation at the interface and poor charge transport. To investigate the effects of B doping on the electron trap state density in the TiO2 films, we conducted cyclic voltammetry (CV) analysis in LiClO4 (0.1 m) solution. Under negative potential (forward bias), the quasi Ef of TiO2 moved toward to Ec, leading to a capacitive current. Ideally, electron injection will commence once the quasi Ef approaches the Ec minimum. However, electrons originating from surface states level below the Ec are injected under forward bias, producing a gradual onset of the capacitive current. Under reverse bias, the capacitive current reached zero at the more positive potentials, implying recovery of the injected electrons and regeneration of most surface states. To sum up, the current peak is related to the occupancy of the surface states in the bandgap.2 sample decreased compared to that of TiO2. This further confirms the reduced electron traps in the presence of B doping.To further understand the surface states of TiO2, resulting partially from the reduction of Ti4+ to Ti3+ . As a result, the inferred Ti3+ defects can induce a shallow energy level below the Ec and act as electron trap sites, which traps electrons and depresses electron transport and conductivity of TiO2. Based on the above measurements and discussion, we show the trap passivation mechanism in Figure 3+ by a species with valency +3.[qv: 11a] Therefore, this substitution by B effectively passivates oxygen vacancy defects in a TiO2 ETL. This conclusion is further supported by the electron paramagnetic resonance (EPR) results as indicated in Figure g = 2.003) is significantly weakened.There is broad consensus that oxygen vacancies or titanium interstitials are the predominant nonstoichiometry defects in TiO2.32 film with reduced defects, the ETL/perovskite interface can be expected to have higher binding energy, that is, be more ideal. To confirm this, we first used DFT to examine interfacial binding at the ETL/perovskite interface  films where we observe a more strongly quenched PL in the B‐TiO2/perovskite stack. This result suggests that the electron extraction from the perovskite was more efficient for B‐TiO2, consistent with the high electron mobility of B‐TiO2 film and improved interfacial binding at the B‐TiO2/perovskite interface. We further investigate the modulation of charge transfer through time‐resolved photoluminescence decay kinetics  can be correlated to the radiative recombination of free carriers due to bulk traps in the perovskite, and the fast decay component (τ1) is primarily dictated by the quenching of carriers at the interface, in the present case involving the ETL/perovskite interface.1) decreases from 8.2 to 4.7 ns for perovskite cast on B‐TiO2 film, indicative of a more rapid interfacial charge transfer in the B‐TiO2/perovskite stack. This result originates from the enhanced electron mobility and improved interface. Meanwhile, the slow decay component (τ2) decreases from 116.8 to 27.0 ns, which may result from a reduced Shockley–Read–Hall recombination via bulk traps in the perovskite film, given the improved film quality with B‐TiO2 ETL .2/perovskite sample, the calculated amplitude of τ2 is decreased from 72.0% to 64.6%, while the calculated amplitude of τ1 is increased from 28.0% to 35.4%, resulting in amplitude average lifetime of 86.4 and 19.1 ns for TiO2 and B‐TiO2 based stack, respectively. Overall, these results further indicate that the interfacial binding and electron extraction are enhanced at B‐TiO2/perovskite interface. In addition, perovskite deposited on quartz substrate shows high PL intensity and long carrier lifetime, indicating the good quality of our perovskite film.Steady‐state photoluminescence measurement was performed to study the charge transfer kinetics between perovskite and ETLs. Figure mponent τ is prima2.42 would make a better ETL in the PSCs than the TiO2. Mesoscopic‐type PSCs were therefore fabricated based on different ETLs with the device structure shown in Figure3 was employed as the active absorber.With the superior optoelectronic properties discussed above, it is expected that the B‐TiOJ–V characteristics of mesoscopic‐type PSCs using different ETLs, with the photovoltaic parameters, including Jsc, Voc, FF, and PCE summarized in Table2 produces a calculated PCE of 19.06% with Jsc = 22.96 mA cm−2, Voc = 1.08 V, and FF = 76.60%. As expected, when the B‐TiO2 is employed as ETL, the Jsc, Voc, and FF are increased to 23.71 mA cm−2, 1.10 V, and 78.60%, yielding a PCE up to 20.51%. The statistical consistency of the performance enhancement was demonstrated by comparing the performance of 24 devices (24 devices each for pure TiO2 and B‐TiO2), and the results are shown in Figure S9 in the Supporting Information. The low device performance for the TiO2 based PSCs is caused by relatively small Jsc and FF, which is related to low electron mobility and higher resistance, and the low Voc results from the carrier recombination at the ETL/perovskite interface (see later discussion). By comparison, the mesoscopic‐type PSCs with B‐TiO2 ETLs exhibit better performance. The higher Jsc and FF are attributed to the high electron mobility that promotes effective electron extraction, and the larger Voc is due to the stronger bonds at the B‐TiO2/perovskite interface that results in reduced interfacial recombination. Figure S10 in the Supporting Information shows the incident photon to current conversion efficiency (IPCE) of the PSCs based on different ETLs. Integration of the IPCE spectra over the solar emission yields AM 1.5 photocurrent of 21.93 and 22.60 mA cm−2 for the device using TiO2 and B‐TiO2, respectively, which agrees well with the Jsc values obtained from the J–V curves within ≈5% deviation. It can be seen that the integrated current density is lower than what is measured from the J–V curves with and without boron doping. This trend has been systematically observed also from other groups, and it can in part be attributed to a small spectral mismatch between the solar simulator and the standard AM 1.5G emission.Figure 2 and B‐TiO2 was measured when the devices were biased at 0.89 and 0.92 V, respectively. Figure J–V plots. The PCEs of the champion devices using the TiO2 and B‐TiO2 ETLs stabilize at 18.57% and 20.26% with photocurrent densities of 20.86 and 22.02 mA cm−2, respectively, very close to the values measured from the J–V curves. As shown in Figure 2 based device rose quickly to maximum steady‐state photocurrent. However, photocurrent from the TiO2 based cell needed more than 2 s to fully saturate. It indicates that it takes longer for the TiO2 based device to reach steady state, which is consistent with the larger hysteresis of TiO2 based devices . According to the equivalent circuit model, the current density J is described asJph is the photocurrent density, J0 is the saturation current density at reverse bias, q is the elementary charge, m is the ideality factor, kB is the Boltzmann's constant, T is temperature in degrees Kelvin, Rs is the series resistance, and Rsh is the shunt resistance. When Rs is far less than Rsh, Equation To gain further insight into the photovoltaic characteristics of TiOV/dJ versus (Jsc − J)−1 and the linear fitting curves for TiO2 and B‐TiO2 based devices according to Equation m can be derived from the linear fitting results. Compared to the TiO2 based device, the value of m decreased from 2.30 to 1.85. The ideality factor m depends on carrier diffusion and recombination processes.[qv: 30a] The smaller the value of m, the less recombination is caused by traps in the device. Thus, a smaller m perhaps means that the carrier recombination in B‐TiO2 based device is reduced because of the reduction of trap states in the B‐TiO2 films.Figure 2 and B‐TiO2 ETLs measured at 0.8 V bias under illumination, with the equivalent circuit shown in Figure Rtra (R1) and the low‐frequency component is for the recombination resistance Rrec (R2).Rtra is the perovskite/ETL interface. When we perform numerical fitting, the device with B‐TiO2 ETL shows smaller Rtra of 6.45 Ω and the larger Rrec of 231.4 Ω compared to that of TiO2 based device (Rtra of 8.33 Ω and Rrec of 147.2 Ω). The smaller Rtra is beneficial for electron extraction, and the large Rrec effectively resists charge recombination, which is in agreement with the observations discussed above.The electrical impedance spectroscopy (EIS) measurements were conducted to extract transfer resistance in the devices. Figure 2 ETL, we studied the density of trap states at the perovskite/ETL interface by performing thermal admittance spectroscopy on complete devices. By interrogating the frequency dependent capacitance via TAS measurements we can determine the energetic profile of trap density of state (DOS). As shown in Figure 2 based device shows more than one order of magnitude lower trap density compared to the control device, confirming that the electronic contact between perovskite and TiO2 is enhanced through boron doping.To further investigate the performance improvement by using B‐TiO2.5FigureJ–V curves measured under both reverse‐ and forward‐scan directions with the related photovoltaic parameters listed in Table 2 has almost identical J–V curves with negligible hysteresis, even when it is measured using different scan rates from 1 to 500 mV s−1 . To quantify the hysteresis effect, hysteric index was given byreverse and PCEforward represent PCE for the reverse (from open‐circuit to short‐circuit under the forward bias voltage) and forward (from short‐circuit to open‐circuit under the forward bias voltage) scans, respectively. PSCs based on B‐TiO2 ETL show a decreased HI of 0.01 compared to that of TiO2 based analogue at 0.13.Hysteresis and stability are two pivotal characteristics for the PSCs.2/perovskite interface, one of the major causes for hysteresis, is significantly reduced. Furthermore, the electron mobility of the TiO2 ETL is only 3.30 × 10−5 cm2 V−1 s−1  HTL.Je) is smaller than the hole flux (Jh), which leads to charge accumulation at the TiO2/perovskite interface, as shown in Figure 2 (1.69 × 10−4 cm2 V−1 s−1) ETL and the larger contact area at the ETL/perovskite interface, the Je and Jh can be balanced when the mesoscopic electron conductor (B‐TiO2) is employed as the ETL  isopropoxide (TTIP)  was added dropwise to a solution consisting of 140 mL of deionized water and 60 mL of ethanol in a 500 mL beaker under vigorous stirring, and the formed precipitate solution was stirred intensely for 4 h after the drip process was finished. After that, the mixture was filtered with a 250 mL G5 sand core funnel, and the precipitate was washed three times with deionized water during the filtration process. The precipitate was then transferred to a Teflon round container  containing 90 g trimethylamine aqueous solution (6.58 wt%), and the solution was stirred for 14 h at a temperature of 80–85 °C to form a translucent sol, after which another 10 g of trimethylamine aqueous solution was added and further stirred for 10 min at room temperature. Subsequently, the resulting sol was transferred into Teflon containers  and placed in sealed metal autoclave vessels, which were then placed in an oven for 10 h at 200 °C. The obtained gel was diluted with ethanol and was evaporated on a rotary evaporator to remove most of the water, and the gel was then purified by rinsing with ethanol three times and centrifuged. For the doped samples, 8.04 g boric acid  was dissolved in a mixed solution of deionized water and ethanol, and stirred for 30 min before the addition of TTIP. TiO2 pastes were then prepared according to previously reported procedures.[qv: 16a]Fabrication of Perovskite Devices: All devices were prepared in ambient air. The TiO2 compact (cp‐TiO2) layer was prepared by a spray pyrolysis method: a solution containing 0.4 mL of acetylacetone and 0.6 mL of TTIP in 7 mL of isopropanol was sprayed on cleaned and patterned FTO  substrates at 460 °C using dry air as a carrier gas. Then, the mesoporous TiO2 (mp‐TiO2) and B‐doped TiO2 films were spin‐coated onto the FTO/cp‐TiO2 substrate using as‐prepared pastes and calcined at 510 °C for 30 min in air to remove organic components. The perovskite layer was deposited from a precursor solution containing 1.2 m PbI2  and 1.2 × 0.98 m CH3NH3I  in a mixed solvent of DMF and DMSO, in which the volume ratio of DMF to DMSO was 1:4. The perovskite precursor solution was spin‐coated on FTO/cp‐TiO2/mp‐TiO2 substrate by a consecutive two‐step spin‐coating program at 1100 and 4200 rpm for 15 and 30 s, respectively. During the second step, 110 µL of chlorobenzene was poured onto the spinning substrate 10 s prior the end of the program. The films were then annealed on a hotplate at 100 °C for 15 min. Once cooled down to room temperature, Spiro‐OMeTAD (Xi'an Polymer Light Technology Corp.) was deposited on top of the perovskite layer by spin coating at 4500 rpm for 20 s. The Spiro‐OMeTAD solution was prepared by dissolving 73.53 mg of Spiro‐OMeTAD in 1 mL chlorobenzene, with the addition of 29.30 µL of 4‐tert‐butylpyridine , 17.22 µL of lithium bis(trifluoromethanesulfonyl)imide  solution (500 mg Li‐TFSI in 1 mL acetonitrile), and 6.78 µL of tris(2‐(1H‐pyrazol‐1‐yl)‐4‐tert‐butylpyridine)‐cobalt(III) tris(bis(trifluoromethylsulfonyl)imide)  solution (400 mg FK209 in 1 mL acetonitrile).Characterization: The morphology measurement was taken by SEM  and TEM . The composition and crystal structure were analyzed by XRD . XPS and UPS measurements were performed on the ESCALAB 250Xi (Thermo Scientific) with different radiation source. The film roughness was collected on an atomic force microscope . The contact angle measurement was carried out by a contact angle meter . The UV–vis absorption spectra were recorded with a spectrophotometer . Steady PL was obtained from a fluorescence spectrometer  with the pump light wavelength of 485 nm and the probe light wavelength of 770 nm.J–V curves were recorded by using a Keithley 2420 source meter under simulated sunlight from Newport solar simulator matching the AM 1.5G standard with an intensity of 100 mW cm−2. A nonreflective metal mask with an aperture area of 0.09 cm2 was used to define the active area of the device. The IPCE curves were confirmed as a function of wavelength from 300 to 800 nm . The IS measurement was carried out on an electrochemical workstation  in the frequency range of 10 mHz to 1 MHz under AM 1.5G condition.The authors declare no conflict of interest.SupplementaryClick here for additional data file.


It is expected that the Cl incorporation in 2D perovskites with layered nature would be much more efficient and thus can greatly alter the morphology, optical properties, phase transition, and charge transport; however, studies on those aspects in 2D perovskites remain elusive up to date. Here, a one‐pot solution method to synthesize the Cl‐doped lead‐free 2D perovskite (BA)2SnI4 with various Cl incorporation concentrations is reported and how the Cl incorporation affects the morphology change, photoluminescence, phase transition, and charge transport is investigated. The Cl element is successfully incorporated into the crystal lattice in the solution‐processed perovskite materials, confirmed by X‐ray photoelectron spectroscopy and energy dispersive X‐ray spectroscopy measurements. The temperature‐dependent photoluminescence studies indicate that the emission properties and phase transition behavior in (BA)2SnI4−xClx can be tuned by varying the Cl incorporation concentration. Electrical measurement suggests that the charge transport behavior can also be greatly altered by the Cl doping concentration and the electrical conductivity can be significantly improved under a higher Cl incorporation concentration.The incorporation of chloride (Cl) into methylammonium lead iodide (MAPbI This has been experimentally demonstrated in 2D perovskites that a high Cl incorporation ratio can be achieved by using vapor phase growth method.It has been demonstrated that chlorine incorporation in the MAPbI2SnIxCl4−x (BA = n‐CH3CH2CH2CH2NH3)2D perovskites and investigate how the Cl incorporation concentration influences the morphology, phase transition, optical and charge transport properties in order to optimize device performance. The morphology of the resultant samples changes from plates to needles with the increase of the Cl incorporation concentration. X‐ray photoelectron spectroscopy (XPS) and energy‐dispersive X‐ray spectroscopy (EDS) reveal the presence of Cl element and the formation of the Sn—Cl bonds in our samples, which has not been observed in solution‐processing 3D perovskites up to date. The temperature‐dependent photoluminescence (PL) and differential scanning calorimeter (DSC) studies suggest that the Cl doping concentration can significantly change structural phase transition temperature. Finally, with the increase of the Cl doping ratio, the electrical conductivity is greatly enhanced and the charge transport changes to variable‐range hopping, indicating that Cl doping introduces the extra crystal structure disorders.Herein, we for the first time report on a one‐pot method to synthesize the lead‐free (BA)2SnI4 crystals with different Cl incorporation concentration were synthesized by using the method similar to previous report except that HI solution was replaced by an excess of aqueous HCl/H3PO2 solution in order to incorporate the Cl element into the (BA)2SnI4 crystals.3PO2 can not only prevent the oxidation of Sn2+ to Sn4+ but also suppress the oxidation of I− to I3− and/or I2. Thus, we properly increase the amount of the H3PO2 solution so as to obtain the (BA)2SnI4−xClx without Sn4+. The ratio of I:Cl in these solutions was gradually increased from 1:0 to 1:8 by adding different amount of BAI solution. It should be noted that the I:Cl ratio stands for the volume ratio of BAI in 57% w/w aqueous HI solution to 37% w/w aqueous HCl solution.(BA)Figure 2SnI4 with the Cl incorporation. As previously reported, the Cl atoms incorporated into the crystal lattice prefer to occupy the apical positions of [PbI6]4− octahedrons but still have a possibility to occupy the equatorial positions of [PbI6]4− octahedrons as well.2SnI4−xClx with different ratio of I:Cl is shown in Figure 2SnI4−xClx evolves from flake‐like to needle‐like with the increase of the Cl concentration in the precursor solution. In addition to the shape evolution, the surface morphology of the resultant (BA)2SnI4−xClx can also be tuned with the Cl concentration. In comparison with the rough surface of the (BA)2SnI4 crystals 2SnI4−xClx needles maintain the layered nature after Cl incorporation 2Se Figure d,e, simin Figure e, which − and I− anions.3−xClx films have the same diffraction peaks as that of pure MAPbI3 films and no Cl—Pb bonds has been observed from XPS spectrum of MAPbI3−xClx films, both of which suggest that the Cl element is not truly incorporated into the crystal lattice of MAPbI3.2SnI4−xClx crystals gradually shift to the smaller angles as the ratio of I to Cl in precursor increases, which agrees with previous reports and suggests2SnI4  peak indicates that two peaks are present with the binding energies of 197.7 and 199.4 eV . Although the I:Cl ratio measured by EDS is not accurate, the trend of the ratio unambiguously indicates the Cl incorporation ratio increases as the increase of the I:Cl ratio in precursor solution. This is the first time that Cl element is successfully doped into the I‐based organic–inorganic halide perovskites by solution method.To further confirm that the Cl element indeed incorporates into the crystal lattice of (BA)4+ within our samples due to the easy oxidation of Sn2+ to Sn4+.2+ without the presence of Sn4+, which might be due to the excess strong reducing H3PO2 solution we used during the growth process . To conclude, all those evidences point to that our one‐pot solution method by using excess H3PO2 solution provides an efficient route to synthesize the Cl‐doped Sn4+‐free tin‐based 2D perovskites for further optoelectronic applications.The XPS spectra were also used to identify whether there is Sn2SnI4−xClx crystals, we have carried out the steady‐state PL studies at room temperature and 77 K. With the increase of the I:Cl ratio, the emission peak continuously shifts from 615 to 630 nm and the full width at half‐maximum (FWHM) simultaneously broadens, which might be due to the Cl incorporation introduced the doping effect, lattice distortion, and disorders.To investigate how the Cl incorporation concentration influences the optical properties of the as‐synthesized (BA)2SnI4−xClx crystals, temperature‐dependent PL studies were further carried out on the exfoliated (BA)2SnI4−xClx microplates. To minimize the degradation of the exfoliated (BA)2SnI4−xClx microplates, we put the samples immediately into the vacuum chamber once they are prepared. Figure 2SnI4 and (BA)2SnI4−xClx with an I:Cl ratio of 1:8 in precursor solution, respectively, excited by a 473 nm continuous wave (CW) laser. For both samples, the emission peak shows a blueshift with the reducing temperature and a sudden peak shift was observed at a certain temperature depending on the cooling and heating cycle, and the Cl incorporation. If we extracted the emission peak for both the cooling and heating cycle, an apparent hysteresis is observed with a temperature span of 20 K for (BA)2SnI4 and 40 K for (BA)2SnI4−xClx with an I:Cl ratio of 1:8, respectively, between the cooling and heating cycle nH2n+1NH3)2MX4 often exhibit a range of temperature‐dependent structural phase transitions, which are associated with changes in the ordering and hydrogen bonding of the organic cations.nH2n+1NH3)2SnI4 compounds affects their optical properties has been investigated before, based on which we are able to attribute the sudden emission peak shift to the orthorhombic‐to‐orthorhombic phase transition with a phase transition temperature of 254.2 K.2NiF4 structure type with a cell setting of orthorhombic (space group: Pbca). The lattice constants for the low‐temperature phase are a = 8.932 Å, b = 26.023 Å, c = 8.408 Å and for the high‐temperature phase are a = 8.814 Å, b = 8.591 Å, c = 27.644 Å.Previous studies reveal that the aliphatic materials .2SnI4−xClx microplates on a 300 nm SiO2/Si substrate with a channel length of 20 µm and SiO2 layer working as dielectric material. Figure 2SnI4−xClx microplates serving as the channel semiconductor. The detailed device fabrication process can be found in the Experimental Section. The inset of Figure 2SnI4−xClx microplate FET device with an I:Cl of 1:8. Unlike the films‐based perovskite FET, in which the presence of grain boundaries and other defects can severely hinder to extract the intrinsic transport parameters, the exfoliated 2D perovskite microplates are single crystals with less defects and thus beneficial for investigating their intrinsic electronic properties.Finally, we have studied how the Cl incorporation influences the charge transport of the as‐synthesized materials, which are indispensable for improving the performance of the optoelectronic devices and designing new device structures.Ids vs Vds) of the 2D perovskite microplate devices with various Cl incorporation concentration at room temperature without applying a gate voltage. To compare among devices, we have normalized the source–drain current by the cross‐section area of the respective microplate. The nonlinearity of the IV curves indicates that the electrical contact is not optimal. Importantly, the source–drain current first decreases and then continuously increases with the increasing Cl incorporation concentration, suggesting that the Cl doping could improve the conductivity of (BA)2SnI4−xClx crystals for a higher Cl incorporation concentration which would be beneficial for the device applications . For all devices, the source–drain current Ids continuously increases with increasing negative gate voltage and decreases with increasing positive gate voltage, indicating a p‐type semiconductor behavior.Figure 2SnI4−xClx microplate devices with various Cl incorporation concentration at a source–drain bias of 2 V at room temperature. All devices exhibit p‐type semiconductor behavior.Figure µ was extracted from the positive sweeping of the transfer characteristics using the following formula: μ=LWCVdsdIdsdVg, where L is the channel length, W is the channel width, C is the gate‐channel capacitance, Vds is the source–drain voltage, Ids is the source–drain current, and Vg is the gate voltage.2SnI4−xClx microplate devices with various Cl incorporation concentration under different source–drain bias. Without Cl incorporation  is linear in T−1/2 from 77 to 180 K  is linear in T−1/2 from 77 to 160 K and linear in T−1/4 from 200 to 294 K. While the T−1/4 dependence behavior above 200 K before phase transition for the devices with the I:Cl ratio of 1:8 might be identified as Mott‐VRH (Mott‐variable range hopping) where transport occurs by thermally assisted hopping between localized electronic states similar to the transport in amorphous Ge, Si2, and FeS2.T−1/2 dependence behavior below 200 K for samples with the I:Cl ratios of both 1:1 and 1:8, however, could be interpreted in a number of ways, particularly conventional Efros–Shklovskii (ES) VRH.For microplates without Cl incorporation, the source–drain current continuously decreases with the increasing temperature for the temperature ranges both after and before phase transition, which originates from the enhanced carrier–phonon interaction with the increase of the temperature Figure g.65 The K Figure h. For th2SnI4−xClx with various Cl incorporation concentration. Cl element has been successfully incorporated into the crystal lattice of (BA)2SnI4, which is the first time to dope Cl into the crystal lattice of the perovskite materials by solution method. It is further demonstrated that the Cl incorporation concentration would alter the morphology of the as‐synthesized crystals, emission peaks, phase transition behavior, and finally charge transport behavior due to the Cl incorporation induced lattice distortion, doping effect, and impurities. Our findings offer an efficient strategy to incorporate the Cl into the lead‐free Sn‐based 2D perovskites with a desired Cl incorporation concentration to improve their performance of electronics and optoelectronics by low‐cost solution method. In addition, this controllable tunability of the Cl incorporation concentration provides not only an extra degree of freedom to modify the properties of the Sn‐based lead‐free 2D perovskites for optotelectronic applications, but also sheds light on the preparation of Sn‐based 2D perovskite thin films with Cl incorporation for the large‐scale solar cells and light‐emitting devices with improved performance.In conclusion, we have synthesized the lead‐free Sn‐based 2D perovskite (BA)Material Preparation: The Sn‐based 2D perovskite bulk crystals with various Cl incorporation concentration were synthesized by using the method similar to previous report except that HI solution was replaced by an excess of aqueous HCl/H3PO2 solution in order to incorporate the Cl element into the (BA)2SnI4 crystals. In details, SnO powder  was dissolved in a mixture of 57% w/w aqueous HI solution (3.5 mL) and 50% aqueous H3PO2 (0.75 mL) by heating to 150 °C under constant magnetic stirring for about 5 min, forming a bright yellow solution. Subsequently, n‐CH3CH2CH2CH2NH3  in 57% w/w aqueous HI solution was then added into the resultant solution and held for another 10 min. The stirring was then discontinued, and the solution was left to cool to room temperature naturally. The as‐grown crystals were finally isolated by suction filtration and thoroughly dried at 60 °C.For the synthesis of the Cl incorporated samples, the exact same procedure was adopted except that 57% w/w aqueous HI solution (3.5 mL) was replaced by 37% w/w aqueous HCl solution (1.5 mL) and BAI solution  for I:Cl = 1:1; HCl (3 mL) and BAI (1 mL) for I:Cl = 1:3; HCl (3 mL) and BAI (0.5 mL) for I:Cl = 1:6; and HCl (4 mL) and BAI (0.5 mL) for I:Cl = 1:8.Material Characterizations: Powder X‐ray diffraction measurements were recorded using a Bruker D2 PHASER . Optical microscopy images were collected by an Olympus BX53M system microscope. The SEM images were acquired on a JEOL 7001F field emission scanning electron microscope equipped with energy‐dispersive X‐ray spectrometer (EDS). DSC measurement was carried out by a PSTA449F3, Netzsch, Germany. XPS was recorded on a Shimadzu/Kratos Axis‐Ultra DLD‐600W instrument equipped with Al/Mg Kα radiation. The photoluminescence measurements were performed in backscattering configuration using a Horiba HR550 system equipped with a 600 g mm−1 excited by a 473 nm solid‐state laser with a power of 0.2 µW. The temperature‐dependent photoluminescence measurements were carried out in a liquid nitrogen continuous flow cryostat (Cryo Industry of America) coupled with a temperature controller. The reflection measurement was carried out in a Horiba HR550 system with a 600 g mm−1 grating illuminated by a 100 W halogen tungsten lamp.Device Fabrication and Electrical Measurement: 5 nm Cr/50 nm Au two‐probe electrodes were defined on a 300 nm SiO2/Si wafer by photolithography and followed by electron beam evaporation and lift‐off process. The 2D perovskite microplates were mechanically exfoliated from their respective bulk crystals by using Scotch tape. The exfoliated 2D perovskite microplates were subsequently transferred onto the predefined electrodes under the aid of microscopy as previously reported.The authors declare no conflict of interest.SupplementaryClick here for additional data file.