Phase Evolution During Perovskite Formation â Insight from Pair

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Phase Evolution During Perovskite Formation – Insight from Pair Distribution Function Analysis Sandy Sanchez, Ullrich Steiner, and Xiao Hua Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.9b00748 • Publication Date (Web): 15 Apr 2019 Downloaded from http://pubs.acs.org on April 15, 2019

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Chemistry of Materials

Phase Evolution During Perovskite Formation – Insight from Pair Distribution Function Analysis Sandy Sanchez1, Ullrich Steiner1, Xiao Hua1* 1: Adolphe Merkle Institute, University of Fribourg, Chemin des Verdiers 4, 1700 Fribourg, Switzerland Abstract: The recent introduction of organometal halide perovskites to solar cells has significantly enhanced the power conversion efficiency of alternative photovoltaic devices, revolutionizing the development of photovoltaic technologies. To produce perovskite thin films with high device performances, various fabrication methodologies have been developed leading to thin films with different surface structures and crystal morphologies. Tremendous efforts have been devoted to characterizing macro- and microscopic structures within these films to better understand the processing-property-performance relationship. However, their atomic structure and its influence on device performance remains poorly understood. To this end, we employed pair distribution function analysis of X-ray total scattering data to obtain crystallographic and compositional information of methylammonium-lead-iodide (MAPbI3) thin films. This analysis revealed a presence of two near-amorphous intermediate phases with local structures that share subtle but significant correlations with the PbI2 precursor and the desired perovskite phase. The structure transformation from these intermediates to the perovskite deviates from the intuitive belief where the molecular cations get inserted between the sheets of layered PbI2 upon the crystallization of perovskite. This knowledge offers

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critical insight into the perovskite formation pathway and reveals an important link between the short-range structure of the thin films and their corresponding device performance. Introduction: Hybrid lead halide perovskite solar cells (PSCs) have been one of the most rapidly advancing photovoltaic technologies in recent years, with a significant enhancement of the power conversion efficiency (PCE) from 3.8 % in 2009 to now 28.0 %1-5. The rapid improvement in device performance includes the development of the high-quality solution-processed perovskite fabrication protocols which enable the preparation of continuous and smooth perovskite thin films with micrometre-sized grains6. These morphological characteristics are believed to enhance the optoelectronic properties of the films, resulting in an improvement of the PCE. As different fabrication methodologies produce thin films with distinct morphologies leading to devices with various PCEs, the study of the processing–property– performance relationship has thus become increasingly important in this field. Due to the advances in the laboratory-based analytical imaging and microscopy techniques, considerable attention has been devoted to the investigation of the macro- and microscopic structure of the thin films and much better knowledge of the perovskite film morphology–optical property relationship has been established. However, the understanding of the atomic structure within these thin films still remains poor, mainly due to the practical challenges involved in the local structure study of these perovskite thin films7 and the lack of routine characterisation tools. A recent study7 of methylammonium lead iodide (MAPbI3) atomic structure using pair distribution function (PDF) via the advanced X-ray total scattering technique revealed that thin films prepared by the so-called antisolvent (AS) method on top of mesoporous TiO2 layers are composed of a bulk and an amorphous component with the latter contributing up to 70% by weight to the film composition. The presence of this highly disordered perovskite has a large influence on the optical properties of the films, resulting in an optical blue shift and


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photoluminescence. These results imply that, in addition to the macroscopic morphological property of the perovskite material, its atomic structure plays an important role in modifying the device performance. The specific motivation to investigate the perovskite atomic structure stems from an observation made during the development of a scalable method to fabricate high-performance perovskite thin films. This method utilises flash infrared annealing (FIRA) which yields crystals with large lateral domain sizes, leading to devices with state-of-the-art efficiencies8, 9. The crystal morphology within FIRA-annealed thin films depends on the duration of IR irradiation. For annealing times between 1.5 – 3.0 s, however, almost undistinguishable film morphologies and powder X-ray diffraction (XRD) patterns were obtained, suggesting very similar structural properties of these thin films on both a macro- and microscopic level. Nevertheless, the photovoltaic performance of devices made from these films differed, implying that the reason behind the variation in device performance might be related to differences in the atomic structure of the perovskite material. In order to acquire a better fundamental understanding of the processing–property–performance relationship of FIRAannealed perovskite films, and to eventually achieve a better control over device fabrication, it is necessary to study the local structure of these thin films. As a tremendously useful tool to characterise nanostructure and atomic distortion, PDF has been proved very powerful to study the short-range structure properties of the battery electrode materials10-12. In this study, we take advantage of its analytical precision on an atomic level to investigate the FIRA-annealed perovskite thin films prepared under various conditions to improve our understanding of the processing-property relationship of this fabrication technique. We have particularly studied two series of thin films prepared either directly on bare fluorine-doped tin oxide (FTO) electrodes or on FTO substrates covered by a mesoporous TiO2 overlayer (denoted as “Meso”) since an earlier PDF study7 showed the



presence of nanostructured TiO2 induced the formation of a relatively large amorphous perovskite component. We also performed the same analysis on the AS-deposited thin films as a comparative reference to establish a correlation between the perovskite atomic structure and the performance of the resulting devices. Results and Discussion: Surface morphology SEM images of the perovskite films deposited directly onto FTO substrates (Fig. 1a) show similar surface structures after FIRA annealing for 1.7, 2.0 and 3.0s. These samples exhibit larger lateral crystal domains compared to the film prepared using the AS method, which is consistent with what was previously reported.8 When annealed on mesoporous TiO2 substrates, only minor changes in the surface morphologies of the AS and FIRA 2.0s perovskite films are seen, but the lateral perovskite crystal sizes of the FIRA 1.7s and 3.0s samples (Fig. 1b) are noticeable larger (~ 100 μm) compared to the films annealed on bare FTO substrates (~ 50 μm). Since the lateral crystal sizes are reasonably monodisperse in all FIRA-annealed films, crystals are nucleated at a well-defined time during the annealing protocol, followed by crystal growth. Differences in the domain size are therefore indicative of either different nucleation properties of the two substrates or different temporal temperature variations within the films during annealing, caused by difference in the heat transfer from the film to the glass substrate across the mesoporous TiO2 layer, or the lack thereof. Evaluation of perovskite crystal structure by XRD The long-range crystal structure of the perovskite films can be evaluated from the powder XRD data. An initial visual comparison across the FTO series (Fig. 2a) exhibits similar XRD patterns for all samples apart from a pronounced reflection observed at a 2θ angle of about


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25.3º in the FTO FIRA 1.4s sample, which is much reduced for longer annealing times. In the Meso series (Fig. 2b), the AS and FIRA 2.0s XRD patterns are similar and comparable to the corresponding FTO samples. A noticeable difference is seen for the Meso FIRA 1.7s and 3.0s patterns, exhibiting a broad feature centred around the 2θ angle of 25º. Since the patterns are corrected for the scattering background arising from the glass capillary, this hump feature should result from an internal component of these two Meso samples. It does not seem to stem from the TiO2 layers since the XRD pattern of the meso-TiO2 reference (Fig. S1) shows welldefined Bragg peaks. Note that this broad feature is accompanied by two discernible Bragg reflections at 2θ angles of 26.5º and 33.8º, which can be ascribed to the (fluorinated) SnO2 (space group P42/mnm) of the FTO glass. There is hence a high probability that this scattering background results from a glassy component within the FTO layer that may have been introduced into the material during the sample preparation. A preliminary XRD refinement was performed on the FTO AS data (Fig. S2) since our previous analysis suggested that it is composed of a pure tetragonal perovskite MAPbI3 (I4/mcm) phase8. Despite some apparent intensity discrepancies, this result shows no indication of a secondary crystalline phase. Therefore, we used the FTO AS pattern as a reference when examining all the other FTO (Fig. 2a) and Meso (Fig. 2b) XRD patterns. A list of all identified phases in this section with their observed Bragg peaks is summarised in table S1. Our analysis revealed two additional phases. New reflection peaks at 2θ angles of 11.8º, 21.7º, and 25.3º are discernible in the Meso FIRA samples with short annealing times (i.e., 1.7 and 2.0 s) and all the FTO samples. They are particularly pronounced in the FTO FIRA1.4s sample. A second phase with peaks at 12.7º and 25.7º only emerges when the annealing time was longer than 2.0 s (i.e., 3.0 and 4.0 s). These observations suggest a sequential formation of two new components, possibly via different mechanisms. Judging by their peak positions, these two phases were tentatively indexed to the trigonal PbI2 (P-3m1),



hereafter denoted as PbI2 (I) and PbI2 (II). While their (0 1 0) reflections seem to be superimposed at 2θ = 21.7º, implying of a similar b-lattice, the (0 0 1) and (0 1 1) peaks of PbI2 (I) are located at lower angles (11.8º and 25.3º respectively), suggesting an expanded clattice or d-spacing between the ab-planes (≈ 7.5 Å) comparing to that of PbI2 (II) (≈ 7.0 Å). The presence of trigonal PbI2 (II) (P-3m1) has previously been confirmed in a perovskite degradation study13, revealing a MAPbI3 decomposition. A phase similar to PbI2 (I) was also observed in several earlier XRD studies14-16 and was attributed to an intermediate formed upon phase transition from PbI2 to MAPbI3. However, the detailed structure of this PbI2 (I) remains unclear; it is believed to be a derivative of the trigonal PbI2 with an expanded clattice16 caused by the intercalation of organic components (i.e., MA, DMF, or DMSO), while others think this phase reflects a deviation from its original trigonal symmetry15. We further performed additional diffraction measurements on a selection of samples using a laboratory X-ray source to verify our observation. The result (Fig. S3) also show a sequential appearance of the PbI2 (I) and (II) phases upon annealing which is unsurprisingly consistent with our synchrotron measurements. Investigation of short-range structure by PDF We proceed with the PDF analysis to gain a better understanding of the local structure of each phase. The PDF patterns from the identically prepared samples on the two surfaces (FTO and Meso) are overlaid in Fig. 3a. Visual comparisons between the FTO and Meso data of the AS and FIRA 2.0s samples shows nearly undistinguishable patterns, suggesting the substrate type does not place a noticeable influence the perovskite phase composition within the thin film nor its short-/medium-range crystal structure. However, the Meso FIRA 1.7s and FIRA 3.0s samples show much weakened PDF intensities compared to the corresponding FTO samples, accompanied by a very intense peak at about 1.6 Å, suggesting a dominant presence of another component. This phase was also seen in most of the other samples with, however,


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much smaller contributions. Surprisingly, the atom-pair distance of this peak does not correspond to any probable bond length within the precursor MAI or the perovskite MAPbI3 structures (see Table S2). It matches however perfectly with the H-I and Si-O distances in the HI (gas) and glass phases, respectively. As the presence of FTO glass was earlier suspected to contribute to the XRD scattering background, the observation of this 1.6 Å PDF peak confirms this speculation. Nevertheless, the presence of certain number of H-I molecules in the system cannot be excluded, although their release into atmosphere is expected17. For a quantitative evaluation, we have carefully studied these PDF data using the least-square refinement. Prior to a systematic investigation, it is important to verify the phase identification established from the XRD analysis. To this end, a series of preliminary refinements were performed on the PDF data of the FTO FIRA 2.0s sample within a short r-range of 2.7 – 10.0 Å (Fig. 3b-d). The initial refinement was performed against the tetragonal MAPbI3 structure (I4/mcm) alone using only isotropic atomic displacement factors (ADPs) for all atom types7. It led to a difference pattern showing multiple peaks above the noise level at about 3.2, 4.6 and 6.3 Å (Fig. 3b). Another refinement including anisotropic ADPs for iodide was also attempted since a previous PDF study of cubic MAPbBr3 and MAPbCl3 reported significantly different anisotropic ADPs for their halides. However, despite slightly improved agreement factor Rw18, these peaks remained unmodeled suggesting they correspond to atom pairs that are not included in the refinement. It indicates the presence of at least one additional phase that is the intermediate PbI2 (I) observed in the corresponding XRD data. Therefore, our next attempt included a trigonal PbI2 (P-3m1) in the structure model. Note that to reduce the number of variables and to minimize interference from correlated parameters, further analyses are all based on isotropic ADPs. The refinement gave rise to a refined PbI2 c-lattice of 7.44 Å, consistent with the value deduced earlier based on the position of the (0 0 1) Bragg reflection. Although this model delivered a better fit to the experimental data (Fig. 3c) reflected by an



improved Rw, the difference pattern still largely resembles that of MAPbI3 alone (Fig. 3b) where the peak at 6.3 Å and the doublet at 3.2 Å remain prominent. This observation implies the local structure of the additional phase does not agree with trigonal PbI2. Other reported intermediates, such as PbI2‧DMF15, (MA)x(DMF)yPbzIx+2z19 and PbI2‧(DMSO)x20, were also tested. However, their short-range structures were also inconsistent with the PDF. In addition, their simulated diffraction patterns disagree with the XRD of the intermediates observed in a previous in situ diffraction study15. Note that the PDF peak of the tetragonal MAPbI3 at 6.3 Å (red pattern in Fig. 3b) almost entirely stems from the nearest Pb-Pb distance between corner-sharing [PbI6] octahedra. This atom pair is however absent in the structure of PbI2 (blue line in Fig. 3c) and other reported (DMF-containing) intermediates due to their edge-sharing [PbI6] linkage (Fig. 3e and 5a). In addition, the perfect [PbI6] octahedron in PbI2 contains six Pb-I bonds with identical lengths of 3.2 Å, which does not lead to a PDF doublet. These observations indicate this intermediate structure may consist of distorted [PbI6] octahedra that are arranged by sharing vertices. We therefore proceeded with the third refinement (Fig. 3d) against a low-dimensional-networked metal halide perovskite21 RxPbIy (with R being organic components such as MA, DMF and/or DMSO) which meets the aforementioned topological requirement. Among the reported RxPbIy phases, two structures of a layered R2PbI4 (R = CnH2n+1NH3 (n = 4, 5, 6)22) were tested including a monoclinic (P121/a1) and an orthorhombic (Pbca) cell. As the monoclinic cell has a β-lattice constant (Fig. 3f) that deviates from a right angle, it was unsurprisingly associated with a better structural flexibility and thus an improved Rw and difference pattern in the refinement (Fig. S4). It is important to note that since the nature and atomic position of the intercalated molecules (R) are undetermined and these molecules are composed of light elements whose PDF contributions are largely negligible in the X-ray data, they are therefore discarded in the refinement. Moreover, in the absence of the crystallographic knowledge of


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this phase, the occupancy of I was tentatively adjusted to 0.75 to ensure a reasonable Pb : I ratio (= 1 : 3) and charge balance. The derived RxPbIy structure from the refinement contains the signature atom pairs mentioned earlier (labelled in Fig. 3f), whose PDF contributions (blue pattern in Fig. 3d) show a remarkable resemblance to the difference PDF (Fig. 3b) and result in a much improved Rw. It is thus evident that this monoclinic RxPbIy offers a better description of the short-range structure of the intermediate which was earlier assigned to PbI2 (I). To verify the presence of this monoclinic phase in other samples, we have additionally performed two series of refinements across all samples against the trigonal PbI2 (Fig. S5a) and monoclinic RxPbIy (Fig. S5b). The improvement from the refinements against RxPbIy within the short r-range is evident in both FTO and Meso series, implying a presence of its shortrange structure at all stages during the sample preparation, even for the samples prepared using the conventional AS method. The PDF contribution of each component derived from these refinements are shown in Fig. 4 for several selected samples (with the remainder shown in Fig. S6). A PbI2 phase with a structure coherence length larger than that of RxPbIy was seen in all FIRA 3.0s and 4.0s samples, consistent with our earlier observation of PbI2 (II) in the XRD data from the same samples. Note that despite the apparent differences between the FTO and Meso PDF profiles of the FIRA 1.7s and 3.0s samples due to an incoherent presence of the FTO components (including a crystalline SnO2 and an amorphous glass phase), they can be both modeled by the refinements. However, a large amount of a low crystallinity glass phase leads to a strongly biased short-range contribution in the PDF, introducing additional challenges for a quantitative analysis of other phases. Given this analytical difficulty, the refinement yielded a satisfactory fit and revealed no noticeable structure differences concerning the perovskite and intermediate phases of the Meso and FTO samples; a similar observation can also be made for the AS and FIRA 2.0s data. These indicate the presence of



mesoporous TiO2 induce little change in the atomic structure of the thin film, despite their somewhat different crystal morphology on a macroscopic scale (Fig. 1), which may arise from different nucleation and crystallization kinetics. Since no noticeable change in the short-range structure was seen between the FTO and Meso samples, they are jointly considered in the subsequent discussion. However, to minimize the interference from the incoherent presence of TiO2 and FTO, we therefore focus on the FTO samples (with the addition of the Meso FIRA 4.0s sample). Structure of the intermediate phases Based on the evolution of the lattice parameters of the intermediate RxPbIy phase derived from the refinement (Fig. S7 and Table S3), we notice considerable crystallographic differences between the RxPbIy structures in various FIRA samples despite no change in the structure space group. These differences are particularly pronounced between the 1.4 and 1.7 s samples, represented by a sudden increase of the b- and c-lattice with a sharp decrease of the a-lattice and β-constant. The increase in the c-lattice along with the reduction in β-constant further results in a sharp increase of the d-spacing between the (0 0 1) planes (Fig. 5e). As the annealing progresses (> 1.7 s), the evolutions of these structure parameters either follow a reverse trend or become less drastic. These observations indicate the RxPbIy structure at 1.4 s is notably different from the other samples. To investigate the similarities and differences between the samples, as well as their crystallographic correlation to the PbI2 precursor and the perovskite, the crystal structures of the relevant phases are visualized in Fig. 5 (and S8). Note that in the following discussion where specific structure differences are explored, the two RxPbIy intermediates are denoted as RxPbIy (α) and (β), corresponding to RxPbIy phases from the FIRA 1.4 s and the FIRA 1.7s – 4.0 s samples, respectively. Particularly, the RxPbIy (FIRA 2.0s) structure is chosen to represent the average structure of the RxPbIy (β) in Fig. 5 (and S8).


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Similar to the layered PbI2 (Fig. 5a), the RxPbIy (P121/a1) also adopts a stacked spatial configuration, where the interlayers are in parallel with the (0 0 1)/ab-planes (Fig. 5b and c). However, they are composed of corner-sharing [PbI6] octahedra instead of an edge-sharing linkage seen in PbI2. Unlike the perfect [PbI6] octahedra in PbI2, the ones in RxPbIy are distorted each consisting of Pb (Wyckoff position: 2a) at the octahedral center with two I1 (4e) and four I2 (4e) atoms at the terminal and bridging positions, respectively. Comparing RxPbIy (α) with (β) (Fig. 5b cf. c), noticeable changes can be observed, mainly in the tilt of the [PbI6] octahedra, leading to a reduced β-constant and an increased d(0

0 1).

This octahedra tilt

associated with the reduced β-constant signifies a topotactic transformation from the layered building block to the final 3D network in the perovskite (Fig. 5d). The progressive change of the d(0 0 1) in RxPbIy implies a modification of the interaction between the interlayer molecular components and the [PbI6] sheet. Although experimental limitations do not allow us to determine the nature of the molecules and their atomic positions, based on earlier studies22, 23, we would expect these molecules (mainly MA) to align in a similar way to those in R2PbI4, forming hydrogen bonds between the ammonium protons and I atoms22. These interactions, despite weak compared to larger organic cations (formamidinium, guanidinium)24, along with the van der Waals forces between the [PbI6] sheets25, result in crystal cohesion along the caxis22 to stabilize the structure. Note that it is difficult to obtain a reliable iodide occupancy in RxPbIy through PDF refinement because of the minor presence of this phase and the compositional and structural complexity of the composite. Hence, the occupancy for both I1 and I2 was initially set off-stoichiometric during the refinement to ensure a sensible (1:3) Pb:I ratio. Interestingly, the perovskite octahedron also contains two I1 (4a) and four I2 (8h) atoms which align perpendicular and along the (0 0 1) planes, respectively. Despite a structure symmetry change from RxPbIy to perovskite, the fact that their [PbI6] units both share two I1 and four I2 atoms indicates that structure coherence is retained upon the phase transition.



Based on this coherence and the multiplicity of each I site, we deduce that those fullyoccupied I1 (4a) and I2 (8h) in perovskite might be translated from half-filled I1 (4e) and fully-occupied I2 (4e) in RxPbIy. This scenario may also offer a cogent explanation to the role of I1 in the transition from a terminal (in RxPbIy) to a bridging (in perovskite) iodide during the topotactic 2D-to-3D conversion. While the spatial transformation from the RxPbIy to MAPbI3 seems evident, there is no clear connection of the arrangement of their [PbI6] units to the PbI2 precursor. In order to unravel a potential structural correlation, it is necessary to convert the PbI2 crystal structure into a 2 × 2 × 1 supercell (Fig. 5a) and to highlight the significantly distorted empty octahedra sites within the interlayers of RxPbIy (α) (Fig. 5b). The link between the two structures thus becomes clearer. If the octahedra in the PbI2 (0 0 1) layer are partially occupied, this configuration seems to resemble the (0 1 0) layer in the RxPbIy (α) structure. Similarly, a selected occupation of the octahedra sites in the PbI2 (0 1 0) layer would resonate with the octahedra arrangement along the RxPbIy (0 0 1) plane. We also noticed that the initial expansion of d(0 0 1) in RxPbIy (α), presumably induced by molecule insertion, appears as an elongation of the interlayer octahedra along the c-axis which leads to an enlarged volume about twice of their original size; despite this notable distortion, the β-angle (~ 116˚) of the RxPbIy (α) unit cell, however, largely retains the original γ-constant (120˚) of the trigonal PbI2 cell. These two factors concurrently result in a d(0 1 3) in RxPbIy (α) that is comparable to the d(0 1 1) in PbI2 (marked in Fig. 6), and consequently in similar peak positions (~ 25˚ 2θ) of their Bragg reflections (Fig. S9). In addition, the particular cell geometry of RxPbIy (α) also yielded an almost identical value of d(2 1 −3) and d(0 1 3). These two reflections thus superimpose, leading to an apparent single peak with a very strong intensity, which rationalizes the observation of a very prominent diffraction peak at 2θ = 25.3˚ in the FIRA1.4 s XRD data (Fig. 2a).


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A very different cell geometry of the (β) structure, despite the same space group, causes substantial shifts of the Bragg peaks, and therefore considerably different diffraction patterns (Fig. S9). While the intense reflections from the (α) phase are centered within the 2θ range of 20 ~ 35˚, the ones for the (β) phase are at lower 2θ angles, between 10 ~ 20˚. These main reflections of both RxPbIy (α) and (β) structures resonate reasonably well with the experimental data of the MAPbI3 precursor and intermediate reported previously in an in situ XRD study15 (Fig. S10), albeit with some discrepancies in their relative peak intensities and positions. Note that since the structure coherence lengths in both structures are very small (~ 20 Å) (Fig. 6b) relative to their unit cell dimensions and their PDFs are always dominated by the atom-pair contribution from a very crystalline perovskite component, it is thus challenging to derive reliable long-range structure solutions of these intermediates solely from the PDF analysis. We do not exclude possibilities of a different space group/super cell as well as any intricate structure features that are common to the layered structures, e.g., stacking defects. That being said, the corner-sharing [PbI6] 2D-framework revealed by our PDF analysis offers critical insight into the full structure derivation of these intermediates. Mechanistic discussion The evolution of the phase fraction (wt %) of each component derived from the PDF refinement as a function of annealing time is shown in Fig. 6a, along with the change of particle sizes of the intermediate RxPbIy and PbI2 phases (Fig. 6b). Based on our structure analysis of RxPbIy, we propose the (α) phase to form at the early stage of annealing. Its local structure is spatially related to PbI2 whose solvated [PbI6] units in the precursor reconnect through vertices forming corner-sharing [PbI6] sheets. These sheets eventually stack, sandwiching MA (and/or other molecules, such as DMF) to form RxPbIy (α). Further incorporation of MA upon annealing triggers a significant tilting of the [PbI6] octahedra giving rise to a conversion from RxPbIy (α) to (β), which eventually results in the



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transformation of the 2D building blocks to form the 3D perovskite network. Within about 2.0 s of FIRA annealing, this 2D-to-3D structure transition reaches its maximum conversion rate producing mixed phases which contain ~ 70 wt% perovskite and ~ 30 wt% RxPbIy (β). Continued exposure to IR (> 2.0 s) however leads to a further increase of temperature, causing the degradation of both perovskite and the intermediate. The final decomposition products are PbI2 for both phases, and no evident formation of any intermediate phase, such as RxPbIy (α), was observed. This implies the degradation mechanism does not follow the reverse route of the perovskite formation via intermediate(s). This is consistent with an earlier in situ XRD study15 and can probably be explained using a layer-by-layer degradation mechanism proposed by in situ ED/TEM studies26, 27. In contrast to the overall agreement on the degradation pathway, the perovskite formation mechanism may vary between different processing techniques. The pathway established from our structural analysis reflects a mechanism of a solution-based synthesis where solvated PbI2 and MA+ coexist in the system. It involves rearrangement of the [PbI6] unit, providing insight into the change of local structure

to

support

previous

mechanistic

study

based

on

long-range

structure

characterization15. This mechanism may differ from the general knowledge of the syntheses based on pre-coated PbI2 films14,

28,

where the molecular components are believed to be

simply inserted within the interlayers of PbI2. In our proposed pathway, the [PbI6] rearrangement results in a re-orientation of the interlayer spacing from that in between the (0 0 1) planes in PbI2 to that in between the (0 0 1) planes in RxPbIy, which are essentially perpendicular to each other (Fig. 5a cf. b). In addition, the transformation of the precursor to MAPbI3 involves (at least) two additional steps via a sequential formation of RxPbIy (α) and (β) intermediates. With better knowledge of the FIRA samples, it allows us to revisit the AS data. Based on the phase composition, the FIRA 2.0s sample shows the highest purity, comparable to that of the


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AS sample, rationalizing their comparable performances in the PSCs8. Their perovskite components show very similar phase compositions despite their different morphology on a macroscopic scale (Fig. 1). The residual intermediate present in these two samples shares similar structure parameters with a near-amorphous coherence length. The same phase compositions and crystallographic properties shared between the AS and FIRA 2.0s data imply a similar perovskite formation mechanism, reflecting the advantage of the FIRA method from the perspective of its scalability and production efficiency. The morphological differences in terms of lateral crystal domain size therefore must stem from a variation in nucleation densities associated with the two preparation techniques8. The shortness of the FIRA heating pulse leads to only a brief supersaturation, limiting the time span for nucleation. This leads to a low nucleation density and thus to large lateral crystal domains, which should be beneficial for PSC performance29. For the intermediate phase, due to its small grain size and low weight percentage, it can be easily overlooked in the structure analysis. Based on our study, comparing between the phase ratio and the morphology of the crystals, it seems that the latter plays a less significant role in controlling the PSC. In a recent structure study of AS-prepared material30, we notice that the difference pattern from the PDF refinement (Fig. S11) contains the unique atom pairs of the RxPbIy structure, indicating its presence. This interesting observation implies the metastability of this component, and more importantly, its common presence in the thin films prepared using either AS or FIRA methods. A better fundamental understanding of this component will help to improve the preparation of perovskite thin films, and to effectively eliminate its presence for further enhancement of the device performance. Conclusion:



We have performed PDF analysis on a series of AS- and FIRA-annealed perovskite thin films deposited onto the FTO substrate with and without an additional layer of mesoporous TiO2. Combining with the XRD data from the X-ray total scattering experiments, we have identified two intermediate phases, RxPbIy (α) and (β), both present in considerable quantities (> 30 wt% within the thin films) exhibiting very short structure coherence lengths (< 2 nm). These intermediates form sequentially upon annealing, adopting a 2D layered stacking arrangement, which seems to be structurally related to the PbI2 precursor and eventually transforms into the 3D perovskite at high temperatures. The presence of the intermediate(s) was observed in all samples, regardless of the thin-film preparation method. Further annealing results in a degradation of the thin film where both the intermediates and the perovskite decompose to PbI2. Adding a layer of mesoporous TiO2 onto the FTO substrate does not seem to modify the perovskite formation thermodynamics under the studied experimental conditions, despite some modification of the crystallization kinetics leading to different lateral domain sizes in the film. Based on our results, the thin films with the best performances, i.e., AS and FIRA 2.0s, are the ones with the lowest RxPbIy concentration, implying the presence of the intermediate and PbI2 (from decomposition) is detrimental for device performance. Therefore, further improvement in film preparation and annealing should target the minimization of the intermediate components and film degradation to increase the overall concentration of the perovskite phase within the thin film. Experiments: Substrate preparation. The FTO coated glass (Pilkington NSG TEC™) was cleaned with Hellmanex soap, followed by 30 min sonication in a Hellmanex 2% water solution, 15 min sonication in IPA, and 5 min of oxygen plasma etching. To prepare TiO2-coated substrates, 30 nm thick TiO2 compact layer was firstly deposited onto the FTO by spray pyrolysis at 450 °C from a precursor solution of titanium diisopropoxide bis (acetylacetonate) in anhydrous


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ethanol and acetylacetonate. After spraying, the FTO substrates were left at 450 °C for 5 min before cooling to room temperature. A mesoporous TiO2 layer was deposited by spin-coating at 4000 rpm for 10 s with a ramp rate of 2000 rpm/s, using a 30-nm particle size TiO2 paste (Dyesol 30 NR-D) diluted in ethanol. After the spin-coating, the substrates were dried at 100 °C for 10 min, followed by annealing under dry air flow on a programmable hotplate (2000 W, Harry Gestigkeit GmbH) at 450 °C for 30 min to crystallize TiO2. Upon cooling down to 150 °C, all the substrates were kept on the fume for the deposition of the perovskite films. Perovskite precursor solutions and thin film preparation. The organic salts were purchased from Dyesol; the lead compounds from TCI. The MAPbI3 perovskite precursor solution contained MAI (1M), PbI2 1M in anhydrous 3 v/v DMF: 1 v/v DMSO. The perovskite antisolvent films were deposited as previously reported

31, 32.

Briefly, the perovskite solution

was spin coated in two steps at 1000 and 6000 rpm for 10 s and 20 s, respectively. During the second step, 100 μL of chlorobenzene was poured onto the spinning substrate 5 s before the end of the program. The substrates were then annealed at 100 °C for 1 h in a nitrogen-filled glove box. After perovskite annealing, the substrates were cooled down within a few minutes. The films made by the FIRA method also followed the procedure reported earlier8. In short, the spin-coating of the perovskite solution involved a single step at 4000 rpm for 10 s. The substrates were then IR irradiated for various time periods (1.4 – 4.0 s) in the FIRA oven and were kept there for 10 additional seconds for cooling before removal. The FIRA processing was carried out in a standard fume hood. For an irradiation period of 1.7, 2.0 and 3.0 s, both FTO and Meso samples were prepared for a comparative study of the influence of the mesoTiO2 layer on the materials’ local structure. For the irradiation of 1.4 and 4.0 s, only FTO or Meso sample was prepared to study structure transition within a larger time (and temperature) window.



Material characterization. Scanning electron microscopy was carried out on a Tescan MIRA 3 LMH with a field emission source operated at an acceleration voltage of 10 kV using an octane-pro EDS detector. X-ray total scattering experiment was performed at beamline I15-1 at Diamond Light Source using an amorphous silicon area detector (PerkinElmer) with an Xray beam of energy 76 keV (λ = 0.1617 Å). Ex situ powder samples were loaded and sealed into borosilicate capillaries (1.0 mm in diameter) and were prepared by carefully scrapping the thin film off the substrates. An empty capillary was measured for background. Data reduction and normalization (Qmax = 24 Å-1) were performed using DAWN33 and PDFgetX234 program, respectively. An additional LaB6 (NIST SRM 660) pattern was collected as a reference to obtain instrumental parameters for XRD Rietveld refinement and the damping factor for PDF refinement. The XRD Rietveld refinement and PDF least-squares refinement were performed using the GSAS II package35 and the PDFgui software36, respectively. Laboratory XRD data were acquired on a STOE STADI P diffractometer using a Cu target (λ = 1.5401 Å) via transmission geometry. Supporting Information XRD pattern of meso-TiO2, XRD refinement of FTO AS sample, laboratory XRD data of selected FIRA samples, PDF refinement methods, results and derived structure parameters for RxPbIy phase, PDF and XRD comparison with other literature. This material is available free of charge via the Internet at http://pubs.acs.org. Acknowledgements We acknowledge the Adolphe Merkle and the Swiss National Science Foundation (Program NRP70 No. 153990) for funding, and Diamond Light Source for time on Beamline I15-1 under Proposal EE17315-1. We thank Prof. Andrew Goodwin, Dr. Hamish Yeung, Dr. Emily Reynolds, Dr. Michael Saliba, Dr. Phoebe Allan and Mr Philippe Holzhey for their time and


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many insightful discussions. We also thank Dr Dean Keeble for his great help in data collection. Conflict of Interest The authors declare no conflict of interest. Corresponding Author Dr. Xiao Hua Address: Adolphe Merkle Institute, University of Fribourg, Chemin des Verdiers 4, 1700 Fribourg, Switzerland Current address: Inorganic Chemistry Laboratory, University of Oxford, South Parks Road, Oxford OX1 3QR, UK Phone: +44(0)1865282774 Email: [email protected]

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8. Sanchez, S.; Hua, X.; Phung, N.; Steiner, U.; Abate, A., Flash Infrared Annealing for Antisolvent-Free Highly Efficient Perovskite Solar Cells. Adv. Energy Mater. 2018, 1702915. 9. Sanchez, S.; Christoph, N.; Grobety, B.; Phung, N.; Steiner, U.; Saliba, M.; Abate, A., Efficient and Stable Inorganic Perovskite Solar Cells Manufactured by Pulsed Flash Infrared Annealing. Adv. Energy Mater. 2018, 1802060. 10. Hua, X.; Liu, Z.; Fischer, M. G.; Borkiewicz, O.; Chupas, P. J.; Chapman, K. W.; Steiner, U.; Bruce, P. G.; Grey, C. P., Lithiation Thermodynamics and Kinetics of the TiO2 (B) Nanoparticles. J. Am. Chem. Soc. 2017, 139, 13330-13341. 11. Pourpoint, F.; Hua, X.; Middlemiss, D. S.; Adamson, P.; Wang, D.; Bruce, P. G.; Grey, C. P., New Insights into the Crystal and Electronic Structures of Li1+xV1–xO2from Solid State NMR, Pair Distribution Function Analyses, and First Principles Calculations. Chem. Mater. 2012, 24, 2880-2893. 12. Hua, X.; Liu, Z.; Bruce, P. G.; Grey, C. P., The Morphology of TiO2 (B) Nanoparticles. J. Am. Chem. Soc. 2015, 137, 13612-13623. 13. Juarez-Perez, E. J.; Ono, L. K.; Maeda, M.; Jiang, Y.; Hawash, Z.; Qi, Y., Photodecomposition and thermal decomposition in methylammonium halide lead perovskites and inferred design principles to increase photovoltaic device stability. J. Mater. Chem. A 2018, 6, 96049612. 14. Pellegrino, G.; D’Angelo, S.; Deretzis, I.; Condorelli, G. G.; Smecca, E.; Malandrino, G.; La Magna, A.; Alberti, A., From PbI2 to MAPbI3 through Layered Intermediates. J. Phys. Chem. C 2016, 120, 19768-19777. 15. Nenon, D. P.; Christians, J. A.; Wheeler, L. M.; Blackburn, J. L.; Sanehira, E. M.; Dou, B.; Olsen, M. L.; Zhu, K.; Berry, J. J.; Luther, J. M., Structural and chemical evolution of methylammonium lead halide perovskites during thermal processing from solution. Energy Environ. Sci. 2016, 9, 2072-2082. 16. Kim, N.-K.; Min, Y. H.; Noh, S.; Cho, E.; Jeong, G.; Joo, M.; Ahn, S.-W.; Lee, J. S.; Kim, S.; Ihm, K.; Ahn, H.; Kang, Y.; Lee, H.-S.; Kim, D., Investigation of Thermally Induced Degradation in CH3NH3PbI3 Perovskite Solar Cells using In-situ Synchrotron Radiation Analysis. Sci. Rep. 2017, 7, 4645. 17. Dualeh, A.; Gao, P.; Seok, S. I.; Nazeeruddin, M. K.; Grätzel, M., Thermal Behavior of Methylammonium Lead-Trihalide Perovskite Photovoltaic Light Harvesters. Chem. Mater. 2014, 26, 6160-6164. 18. Masadeh, A.; Božin, E.; Farrow, C.; Paglia, G.; Juhas, P.; Billinge, S.; Karkamkar, A.; Kanatzidis, M., Quantitative size-dependent structure and strain determination of CdSe nanoparticles using atomic pair distribution function analysis. Phys. Rev. B 2007, 76. 19. Petrov, A. A.; Sokolova, I. P.; Belich, N. A.; Peters, G. S.; Dorovatovskii, P. V.; Zubavichus, Y. V.; Khrustalev, V. N.; Petrov, A. V.; Grätzel, M.; Goodilin, E. A.; Tarasov, A. B., Crystal Structure of DMF-Intermediate Phases Uncovers the Link Between CH3NH3PbI3 Morphology and Precursor Stoichiometry. J. Phys. Chem. C 2017, 121, 20739-20743. 20. Manser, J. S.; Saidaminov, M. I.; Christians, J. A.; Bakr, O. M.; Kamat, P. V., Making and Breaking of Lead Halide Perovskites. Acc. Chem. Res. 2016, 49, 330-338. 21. Saidaminov, M. I.; Mohammed, O. F.; Bakr, O. M., Low-Dimensional-Networked Metal Halide Perovskites: The Next Big Thing. ACS Energy Lett. 2017, 2, 889-896. 22. Billing, D. G.; Lemmerer, A., Synthesis, characterization and phase transitions in the inorganic-organic layered perovskite-type hybrids [(CnH2n + 1NH3)2PbI4], n = 4, 5 and 6. Acta Crystallogr. B 2007, 63, 735-747. 23. Billing, D. G.; Lemmerer, A., Synthesis, characterization and phase transitions of the inorganic–organic layered perovskite-type hybrids [(CnH2n+1NH3)2PbI4] (n = 12, 14, 16 and 18). New J. Chem. 2008, 32, 1736-1746. 24. Svane, K. L.; Forse, A. C.; Grey, C. P.; Kieslich, G.; Cheetham, A. K.; Walsh, A.; Butler, K. T., How Strong Is the Hydrogen Bond in Hybrid Perovskites? J. Phys. Chem. Lett. 2017, 8, 61546159. 25. Tubbs, M. R., The optical properties and chemical decomposition of halides with layer structures. II. defects, chemical decomposition, and photographic phenomena. Phys. Status Solidi B 1975, 67, 11-49.


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26. Fan, Z.; Xiao, H.; Wang, Y.; Zhao, Z.; Lin, Z.; Cheng, H.-C.; Lee, S.-J.; Wang, G.; Feng, Z.; Goddard, W. A.; Huang, Y.; Duan, X., Layer-by-Layer Degradation of Methylammonium Lead Tri-iodide Perovskite Microplates. Joule 2017, 1, 548-562. 27. Kim, T. W.; Shibayama, N.; Cojocaru, L.; Uchida, S.; Kondo, T.; Segawa, H., Real-Time In Situ Observation of Microstructural Change in Organometal Halide Perovskite Induced by Thermal Degradation. Adv. Funct. Mater. 2018, 0, 1804039. 28. Jain, S. M.; Philippe, B.; Johansson, E. M. J.; Park, B.-w.; Rensmo, H.; Edvinsson, T.; Boschloo, G., Vapor phase conversion of PbI2 to CH3NH3PbI3: spectroscopic evidence for formation of an intermediate phase. J. Mater. Chem. A 2016, 4, 2630-2642. 29. Huang, F.; Dkhissi, Y.; Huang, W.; Xiao, M.; Benesperi, I.; Rubanov, S.; Zhu, Y.; Lin, X.; Jiang, L.; Zhou, Y.; Gray-Weale, A.; Etheridge, J.; McNeill, C. R.; Caruso, R. A.; Bach, U.; Spiccia, L.; Cheng, Y.-B., Gas-assisted preparation of lead iodide perovskite films consisting of a monolayer of single crystalline grains for high efficiency planar solar cells. Nano Energy 2014, 10, 1018. 30. Masi, S.; Aiello, F.; Listorti, A.; Balzano, F.; Altamura, D.; Giannini, C.; Caliandro, R.; Uccello-Barretta, G.; Rizzo, A.; Colella, S., Connecting the solution chemistry of PbI2 and MAI: a cyclodextrin-based supramolecular approach to the formation of hybrid halide perovskites. Chem. Sci. 2018, 9, 3200-3208. 31. Ball, J. M.; Lee, M. M.; Hey, A.; Snaith, H. J., Low-temperature processed mesosuperstructured to thin-film perovskite solar cells. Energy Environ. Sci. 2013, 6, 1739-1743. 32. Snaith, H. J.; Abate, A.; Ball, J. M.; Eperon, G. E.; Leijtens, T.; Noel, N. K.; Stranks, S. D.; Wang, J. T.-W.; Wojciechowski, K.; Zhang, W., Anomalous Hysteresis in Perovskite Solar Cells. J. Phys. Chem. Lett. 2014, 5, 1511-1515. 33. Basham, M.; Filik, J.; Wharmby, M. T.; Chang, P. C. Y.; El Kassaby, B.; Gerring, M.; Aishima, J.; Levik, K.; Pulford, B. C. A.; Sikharulidze, I.; Sneddon, D.; Webber, M.; Dhesi, S. S.; Maccherozzi, F.; Svensson, O.; Brockhauser, S.; Náray, G.; Ashton, A. W., Data Analysis WorkbeNch (DAWN). J. Synchrotron Radiat. 2015, 22, 853-858. 34. Qiu, X.; Thompson, J. W.; Billinge, S. L. J., J. Appl. Crystallogr. 2004, 37, 678. 35. Toby, B. H.; Von Dreele, R. B., GSAS-II: the genesis of a modern open-source all purpose crystallography software package. J. Appl. Crystallogr. 2013, 46, 544-549. 36. Farrow, C. L.; Juhas, P.; Liu, J. W.; Bryndin, D.; Božin, E. S.; Bloch, J.; Proffen, T.; Billinge, S. J. L., PDFfit2 and PDFgui: computer programs for studying nanostructure in crystals. J. Phys.: Condens. Matter 2007, 19, 335219.



Figures

Figure 1. SEM images of the a) FTO and b) Meso AS- and FIRA-annealed samples prepared with different IR irradiation times. Yellow dashed rectangles mark the crystal domains in the FIRA samples with defined grain boundaries whose dimensions are representative of the average domain size of each sample.


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Figure 2. XRD patterns of a) FTO and b) Meso samples obtained from the X-ray total scattering experiment. Asterisks mark the main reflections of SnO2 (P42/mnm) from the FTO substrate. The highlighted regions at low 2θ are magnified on the right to better visualize the Bragg reflections of PbI2 (I) and (II) (P-3m1), labelled with red and blue dotted lines, respectively.



Figure 3. a) PDF patterns of the AS and FIRA samples from FTO (black circles) and Meso (solid lines) thin films. The color of the Meso samples follows the color code of Fig. 2. The highlighted peak at 1.6 Å corresponds to Si-O atom pair from the FTO glass. Red triangles mark the distinct atom pairs from SnO2. Short-range PDF refinements (2.7 – 10.0 Å) of the FTO FIRA 2.0s data against model b) MAPbI3, c) MAPbI3 + PbI2, and d) MAPbI3 + RxPbIy where experimental, calculated, simulated (for PbI2 and RxPbIy) and difference patterns are shown in black circles, red, blue and green solid lines, respectively. The discernible PDF peaks in the difference patterns are marked by black triangles. The main atoms pairs in the simulated patterns using e) PbI2 and f) RxPbIy (R molecules omitted) structures are indicated using triangles with various colours, while the corresponding atom pairs in the structures are labelled using the same color code.

Figure 4. Refinements of PDF patterns of selected samples: a) FTO AS, b) FTO FIRA 1.4s, c) FTO FIRA 2.0s and d) Meso FIRA 4.0s, where experimental, calculated and difference pattern are shown as black circles, red and green solid lines, respectively. The PDF contribution from MAPbI3, RxPbIy


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and PbI2 are shown ein blue, orange and purple solid lines, respectively. Note that the y-axis has two different scales.

Figure 5. [PbI6] structures of a) PbI2, b) RxPbIy (α), c) RxPbIy (β) and d) MAPbI3 viewed from different directions. Note that the unit cell lattice vectors are different between PbI2 and the other structures. [PbI6] units are shown as dark grey octahedra where Pb and I atoms are denoted by grey and purple spheres, respectively. R molecules (including MA) are omitted for a clearer view. Unit cells are marked with dotted blue lines. For PbI2, the original unit cell is marked using white dashed lines; the blue dotted lines indicate the supercell. Two different iodine atoms in RxPbIy and MAPbI3 are labelled as I1 and I2. The important crystal planes are marked and labelled with different colors. Note that the light grey sites in a) and b) are empty sites; they are highlighted to show the correlation between the PbI2 and RxPbIy (α) structures. e) The evolution of the d-spacing between the (0 0 1) planes in RxPbIy.

Figure 6. Evolution of the a) weight fraction of each component and b) particle sizes (reflecting the structure coherence lengths) of PbI2 and RxPbIy derived from PDF refinements.



TOC


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Phase Evolution During Perovskite Formation â Insight from Pair

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