Crystalline Intermediates and Their Transformation Kinetics during the

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Crystalline Intermediates and Their Transformation Kinetics during the Formation of Methylammonium Lead Halide Perovskite Thin Films Petr P. Khlyabich, and Yueh-Lin Loo Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.6b04020 • Publication Date (Web): 19 Nov 2016 Downloaded from http://pubs.acs.org on November 22, 2016

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Chemistry of Materials

Crystalline Intermediates and Their Transformation Kinetics during the Formation of Methylammonium Lead Halide Perovskite Thin Films Petr P. Khlyabich1 and Yueh-Lin Loo1,2* 1

Department of Chemical and Biological Engineering, Princeton University,

Princeton, New Jersey 08544, United States 2

Andlinger Center for Energy and the Environment, Princeton University, Princeton, New Jersey

08544, United States

ABSTRACT. The morphology of methylammonium lead halide perovskite thin films significantly affects the performance of optoelectronic devices that comprise them. Using x-ray diffraction studies, we elucidated the mechanisms of thin-film formation to complete the complex picture of structural development of perovskites under inert atmosphere. The presence of excess methylammonium iodide during perovskite crystallization leads to the formation of layered intermediates and low-dimensional perovskites; these intermediates are correlated with the formation of large and continuous grains in the final films. When the precursors are present in stoichiometric equivalence, initial stages of crystallization instead involve the formation of solvates; the fully crystallized films have poor surface coverage and comprise needle-like

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structures. The activation energy of crystallization in films with stoichiometric excess methylammonium iodide is higher than that for films comprising stoichiometrically equivalent precursors; the higher energy barrier is consistent with the need to sublime excess methylammonium iodide during film formation. Replacing lead iodide with lead chloride does not qualitatively alter the crystallization process or the final morphology; it lowers the activation energy for crystallization, presumably because sublimation of methylammonium chloride is less energetic than that of methylammonium iodide. Our study suggests the necessity of layered structures and low-dimensional perovskites for the formation of technologically-relevant continuous thin films.

Scientific interest and research surrounding photovoltaic devices comprising organometal halides (CH3NH3PbI3) that adopt the perovskite structure is currently experiencing an enormous upsurge.1–11 Perovskite thin films possess a unique set of attributes, including opto-electronic properties that are tunable with chemistry and composition,1,5 low exciton binding energy on the order of 1-3 kBT,12–14 large free-charge-carrier diffusion lengths,6,15–17 and high mobilities.6,18–20 Perovskite active layers are solution-processable at or near ambient conditions, not unlike many of the organic and polymer semiconductors, thus preserving the simplicity and low-cost aspects of processing. The high power-conversion efficiencies of devices comprising these active layers stem from high current densities, which can be attributed to strong and uniform absorption in the visible and the near-infrared,21 as well as open-circuit voltages (Voc) often in excess of 1 V,5 given their electronic band gaps. At the time of this writing, efficiencies of such perovskite solar cells have exceeded 18%, with a champion efficiency of 22.1%.22–25


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Despite recent achievements, perovskite photovoltaics suffer a number of drawbacks originating from a lack of understanding of the morphological and structural evolution of the active layer during and after film formation.26 State-of-the-art solar cells demonstrate poor stability, reliability and reproducibility27–29 in large part because the active-layer morphology continues to evolve under illumination when the device is under operation.30–33 This morphological instability often manifests itself in the form of highly hysteretic output characteristics in solar cells.14,34–38 Even though a recent publication reported a >1 cm2 solar cell,39 the majority of the highefficiency devices still have impractically small footprints (< 0.1 cm2) as larger cell areas often incorporate pinholes and defects in the active layers that cause electrical shorts in devices. Our lack of understanding of the film-formation process of these materials is further accentuated by the myriad of processing parameters and conditions, ranging from different lead salt precursors (PbX2) and solvent choice, stoichiometric/non-stoichiometric use of precursors, and postdeposition treatments, that have been employed to access active layers for high performance devices.24,40–44 Empirical studies have demonstrated that the final solid-state morphologies of perovskite thin films are highly dependent on their processing parameters. Most notably, the presence of excess methylammonium iodide (CH3NH3I) in precursor solutions results in continuous and smooth perovskite films with micrometer-sized grains. In contrast, thin films cast from solutions with stoichiometrically equivalent precursors tend to have poor coverage.26 Recent mechanistic studies of CH3NH3PbI3 formation reported the presence of crystalline precursors and structural intermediates at early stages of thin-film formation in the presence of excess methylammonium iodide (CH3NH3I).25,45–51 Films that are processed from solutions having stoichiometrically


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equivalent precursors are instead characterized by the presence of an unidentified crystalline phase.25,48 While these studies begin to shed light on the perovskite film formation, the transformation pathways remain not well understood. Moreover, the kinetics by which different crystalline intermediates transform to the final tetragonal phase of CH3NH3PbI3 is not established. Our study elucidates the kinetics and pathways by which organometal halide perovskites crystallize and establishes strong correlations between different intermediates and the final thin-film morphologies. Results and Discussion Given that both lead iodide (PbI2) and lead chloride (PbCl2) are common precursors for perovskite active layers in photovoltaics, we looked at perovskite thin films derived from both lead salts with CH3NH3I. These films are formed by spin-coating co-solutions of a 1:3 molar ratio of PbX2 and CH3NH3I precursors in N,N-dimethylformamide (DMF). To examine the case in which the precursors are in stoichiometric equivalence, we reduced the fraction of CH3NH3I. We only formed films having a stoichiometric equivalence of PbI2 and CH3NH3I; at stoichiometric equivalence, PbCl2 is not completely soluble in DMF. This observation points to the solvating power of CH3NH3I.


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Figure 1. (a) X-ray diffraction traces illustrating the solid-state structural development of thin films comprising a 1:3 molar ratio of PbI2 and CH3NH3I during thermal annealing at 162°C. (b) High-resolution x-ray diffraction traces illustrating the evolution of the (220) reflection of the tetragonal phase of methylammonium lead iodide for the same thermal annealing condition. Figure 1a shows the structural evolution of perovskite films derived from a 1:3 molar ratio of PbI2 and CH3NH3I upon thermal annealing at 162°C under inert atmosphere. The x-ray diffraction traces were obtained after the thin films had been annealed for a pre-specified time at 162°C and then cooled to room temperature to ensure that we capture the kinetics associated with the formation of the desired tetragonal phase of the perovskite. The x-ray diffraction trace of the as-cast thin film comprises three peaks at 2θ of 9.16°, 18.32° and 27.48°. These x-ray diffraction


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peaks belong to the same family of planes and suggest the presence of a layered structure whose arrangement is periodic in the out-of-plane direction. The diffraction trace of the as-cast film does not include reflections of either of the precursors (PbI2 or CH3NH3I), or CH3NH3PbI3 in its desired tetragonal crystal structure with 2θ at 14.12, 28.44 and 42.30° (Figure S1).52,53 After thermal annealing for 3 minutes at 162°C and quenching to room temperature, the layered structure transforms into the desired tetragonal crystal structure of CH3NH3PbI3. At the early stages of CH3NH3PbI3 crystallization, one additional set of reflections is present in its x-ray diffraction trace. The high-resolution x-ray diffraction trace in Figure 1b reveals a peak at 2θ = 28.24°, as opposed to at 2θ = 28.44° that is assignable to the (220) peak of tetragonal CH3NH3PbI3. This reflection, along with others at 2θ of 14.02° and 42.94°, is slightly shifted compared to the reflections associated with the tetragonal phase, yet they cannot be assigned to the commonly observed cubic phase of CH3NH3PbI3. Moreover, the (111) reflection that is uniquely associated with the cubic phase at 2θ = 24.33° is absent in all the x-ray traces.54 Instead, these reflections indicate the presence of a metastable polymorph of CH3NH3PbI3, referred to as α-CH3NH3PbI3 in the literature.55,56 Increasing annealing time eliminates the presence of this metastable polymorph. The film comprises only tetragonal CH3NH3PbI3 upon complete conversion. The crystallization of CH3NH3PbI3 in the presence of excess CH3NH3I is complex. We observe the formation of a layered structure upon spin-coating that transforms to a coexistence of the desired tetragonal phase of CH3NH3PbI3 with the α-polymorph of CH3NH3PbI3 upon thermal annealing. With extended annealing, the α-polymorph of CH3NH3PbI3 transforms to the tetragonal phase. This transformation pathway is schematized in Figure 4a.


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Figure 2. (a) X-ray diffraction traces illustrating the solid-state structural development of thin films comprising a 1:1 stoichiometric ratio of PbI2 and CH3NH3I during thermal annealing at 74°C. (b) High-resolution x-ray diffraction traces illustrating the evolution of the (220) reflection for the same thermal annealing condition. Next, we examined the mechanism of perovskite thin-film formation of our reference system that comprises a stoichiometric equivalence of PbI2 and CH3NH3I. Comparison with the preceding case sheds light on the role excess CH3NH3I plays on perovskite structural development. Ideally, we should conduct thermal annealing at temperatures relevant to solar-cell fabrication.57 Given the rapid crystallization of perovskites at this precursor molar ratio, however, we had to conduct our experiments at a substantially lower temperature range of 52°C - 74°C. The x-ray diffraction


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trace of the as-cast thin film having a stoichiometric equivalence of precursors comprises three intense peaks at 2θ of 6.48°, 8.00° and 9.48° (Figure 2a). These reflections are consistent with those of a CH3NH3I:PbI2:DMF solvate.58,59 The diffraction trace of the as-cast film does not include reflections of either of the precursors (PbI2 or CH3NH3I), or CH3NH3PbI3 in its desired tetragonal phase (Figure S1).52,53 Upon thermal annealing for 1 minute at 74°C, reflections corresponding to the tetragonal crystal structure of CH3NH3PbI3 emerge alongside peaks associated with the solvate. Increasing the annealing time leads a strengthening of the reflections associated with the desired tetragonal phase of CH3NH3PbI3 (Figure 2b) and a concomitant decrease in the intensity of reflections associated with the solvate. Thus, thermal annealing of the film must trigger gradual evaporation of DMF and subsequent conversion of the solvate to the desired tetragonal phase of CH3NH3PbI3,22,40,60 as illustrated schematically in Figure 4b.


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Figure 3. (a) X-ray diffraction traces illustrating the solid-state structural development of thin films comprising a 1:3 molar ratio of PbCl2 and CH3NH3I during thermal annealing at 132°C. (b) High-resolution x-ray diffraction traces illustrating the evolution of the (220) reflection for the same thermal annealing condition. The use of PbCl2 as an alternative lead-salt precursor has become increasingly popular in recent years.10 The x-ray diffraction traces of fully crystallized films comprising PbCl2 as the precursor, shown in Figure S1, reveal only a single family of out-of-plane reflections. This observation is consistent with prior findings61,62 that replacing PbI2 with PbCl2 leads to more preferentiallyoriented films. The x-ray diffraction trace of the as-cast film having a 1:3 molar ratio of PbCl2 and CH3NH3I (Figure 3a) exhibits reflections at 2θ of 5 - 7°, 11 - 13°54 and 16 - 19°.46 These reflections were previously assigned to low-dimensional perovskites, as well as crystalline CH3NH3PbCl3 with its (100) reflection at 15.60°.53 Importantly, these reflections are distinct from those observed during the crystallization of CH3NH3I and PbI2 in which layered structures are observed (Figure 1a). At early stages of thermal annealing, the intensities of the reflections associated with the low-dimensional perovskites increase. Figure 3b captures the disappearance of the reflection that is associated with crystalline CH3NH3PbCl3 and the simultaneous appearance of reflections corresponding to both α-CH3NH3PbI3 and the desired tetragonal phase of CH3NH3PbI3. The observation of the progressive disappearance of the peak associated with crystalline CH3NH3PbCl3 upon thermal annealing contradicts earlier studies where the (100) reflection of CH3NH3PbCl3 is preserved during thermal annealing and is retained in fullyconverted films.63,64 We speculate that differences in the processing conditions, in particular whether thermal annealing was carried out in air or in nitrogen, strongly affect the lifetime of CH3NH3PbCl3. After thermal annealing for 5 minutes at 132°C, the low-dimensional perovskites


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are converted to a coexistence of its metastable α-polymorph and the desired tetragonal phase of CH3NH3PbI3. Further increasing the annealing time eliminates the α-polymorph, as evidenced by the disappearance of the reflection at 2θ = 28.24° in Figure 3b, leading to films comprising only the desired tetragonal phase of CH3NH3PbI3 (2θ = 28.44°). This transformation process is schematized in Figure 4c. Despite the fact that x-ray diffraction traces of all three crystallized films indicate complete conversion to CH3NH3PbI3 (Figure S1), the morphology of these perovskite films is affected significantly by processing conditions, as demonstrated in the optical micrographs and scanning electron microscopy (SEM)images in Figures S3 and S4. The micrographs of films derived from a stoichiometric excess of CH3NH3I (independent of which lead salt) reveal two-dimensional grains with excellent surface coverage. In particular, the morphology of the perovskite thin films obtained with a 1:3 molar ratio of PbI2 and CH3NH3I, as seen in the optical micrograph in Figure S3b, comprises sub-1 µm grains with uniform surface coverage over large areas. The optical micrograph of the perovskite thin film from a 1:3 molar ratio of PbCl2 and CH3NH3I in Figure S3c reveals dendritic grains, also with continuous coverage of the substrate. In contrast, the micrograph of the film derived from a stoichiometric equivalence of precursors reveals needles, with little coverage between grains. These morphologies are not inconsistent with reports that empirically correlate increasing surface coverage with increasing concentrations of CH3NH3I.26 Perhaps not coincidentally, our x-ray diffraction measurements point out that films with excellent surface coverage are characterized with the presence of layered structures and lowdimensional perovskites in the early stages of film formation. Such intermediates are instead absent during the structural evolution of films whose final morphology comprises onedimensional needles. Given that these low-dimensional perovskites (2D perovskites) and layered


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structures are preferentially oriented with their layers in the plane of the substrate, we speculate that this orientation must promote rapid lateral grain growth, ultimately yielding full surface coverage. It follows that the presence of excess CH3NH3I has to play a role in this growth, likely by plasticizing these films, thereby allowing for lateral diffusion of precursors during the course of solidification. In the absence of excess CH3NH3I, solidification takes place almost immediately upon film casting, with little to no possible rearrangement. This hypothesis is further consistent with the quantification of kinetics of CH3NH3PbI3 formation, as detailed below.

Figure 4. Schematic representation of the structural development during perovskite film formation for (a) PbI2:CH3NH3I = 1:3, (b) PbI2:CH3NH3I = 1:1 and (c) PbCl2:CH3NH3I = 1:3. The heights of these triangles are qualitative proxies for the relative fractions of the crystalline intermediates; positive slopes thus indicate growth in the fraction while negative slopes indicate decreases in the fraction of the crystalline intermediates. The optical micrographs on the right of each process demonstrate the types of mesostructures obtained after crystallization.


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We examined the kinetics to quantify the effects the lead salt and the presence of excess CH3NH3I have on thin-film formation. In all cases, PEDOT:PPS-coated glass was used as the substrate. As a result, the crystallization kinetics obtained in current study may vary from previously reported data where different substrates (e.g., TiO2, ZnO, C60, etc.) were used. We examined the kinetics of perovskite thin-film formation by tracking the evolution of intensity associated with the (220) reflection, as opposed to the (110) reflection, of the tetragonal CH3NH3PbI3 phase as a function of annealing time. We chose to track the (220) reflection because it is better resolved from the second reflection of the metastable polymorph of CH3NH3PbI355,56 compared to their corresponding first reflections. As can be seen in Figure 5, the crystallization of perovskite thin-film formation comprising 1:3 molar ratio of PbI2 and CH3NH3I follows first-order Avrami kinetics. Interpretation of the Avrami exponent (n) in this case is challenging, since crystallization is simultaneous to the chemical reactions that are taking place. Assuming crystallization is the dominant mechanism we are tracking, first-order kinetics can indicate either one-dimensional growth assuming instantaneous nucleation at the start of thermal annealing, or instantaneous growth with nucleation taking place during the course of annealing. As seen in the optical micrograph in Figure S3b, the morphology of the perovskite thin films obtained with a 1:3 molar ratio of precursors comprises sub-1 µm grains, suggesting isotropic growth along its different in-plane crystallographic directions. In the absence of needle-like crystals, we have to rule out onedimensional growth and are left to surmise that nucleation must take place over the course of annealing. Not to be ruled out is also the possibility that chemical reactions dominate; the Avrami exponent here would thus reflect a first-order chemical reaction that is taking place between the precursors. By tracking the evolution of the intensity of the desired (220) reflection


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at different temperatures, and by fitting these data sets to the Arrhenius equation, we found the activation energy for CH3NH3PbI3 formation to be 24 ± 2 kcal/mol. With thin films processed from stoichiometrically equivalent precursors of PbI2 and CH3NH3I, the analysis of perovskite crystallization kinetics reveals first-order (n = 1) Avrami kinetics (Figure S2). Unlike the previous case, the morphology of films that are derived from having a stoichiometric equivalence of precursors reveals needle-like crystals (Figure S3a). We thus believe instantaneous nucleation and one-dimensional growth to be the mechanism responsible for film formation during thermal annealing here. By carrying out thermal annealing experiments at different temperatures (Figure S2), we estimated the activation energy for perovskite crystallization for films comprising 1:1 PbI2 and CH3NH3I to be 4.1 ± 0.2 kcal/mol. This activation energy is only 7 times higher than kT (0.6 kcal/mol) at room temperature. We also note that the activation energy for perovskite crystallization in this case is 6 times smaller than the activation energy for crystallization into its tetragonal phase in the presence of excess CH3NH3I. This comparison reveals that CH3NH3PbI3 is more readily formed with stoichiometric equivalence of precursors. This observation also suggests that the removal of excess CH3NH3I during film formation significantly increases the barrier to perovskite crystallization. Unlike the kinetics of perovskite formation with PbI2, the kinetics of CH3NH3PbI3 thin-film formation with PbCl2 is second order in nature (n = 2), as shown in Figure S2. Assuming instantaneous nucleation, a second-order Avrami exponent indicates two-dimensional growth. The optical micrographs in Figure S3c reveal dendritic crystals that are consistent with this assertion. The estimated activation energy for CH3NH3PbI3 crystallization is 20.4 ± 0.4 kcal/mol. That its activation energy is lower than that for film formation with PbI2 is consistent with the observation that methylammonium chloride (CH3NH3Cl) is more volatile than CH3NH3I.65 Its


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removal is thus less energy-intensive and incurs a smaller barrier compared to removing excess CH3NH3I.

Figure 5. Evolution of the normalized integrated intensities of the (220) reflection during thermal annealing of thin films comprising a 1:3 molar ratio of PbI2 and CH3NH3I at 144°C, 152°C and 162°C. The solid lines are fits to a first-order Avrami equation. Despite differences in crystallization pathways that these films undertake when crystallized in nitrogen and when crystallized in air,25,45,46,66 the x-ray diffraction traces of the fully crystallized CH3NH3PbI3 films are similar and independent of the atmosphere in which solidification takes place. Moreover, the activation energies for the perovskite films formation appear to be comparable whether CH3NH3PbI3 is formed in air45,67 and under inert atmosphere (current study). This similarity in activation energies despite differences in crystallization pathways suggests that the rate-limiting step – whether solidification takes place in air or in nitrogen – is the same. Transitions between intermediate states with lower free energies, which appear to be


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distinct given differences in crystallization pathways, are not accounted for in the quantification of activation energies.68 Conclusions In conclusion, our study provides pertinent information that complements and enhances existing understanding of structural development of perovskite thin films. Our structural studies suggest the criticality of excess CH3NH3I in perovskite crystallization. It induces the formation of layered intermediates and low-dimensional perovskites that we interpret are requisite precursors to large and continuous two-dimensional grains in the final films. Subsequent removal of excess CH3NH3I necessitates higher annealing temperatures and results in larger activation energies, compared to those needed to process films comprising a stoichiometric equivalence of precursors. At stoichiometric equivalence of precursors, the presence of CH3NH3I:PbI2:DMF solvate at early stages of crystallization facilitate the formation of needle-like crystals, ultimately resulting in poor coverage of the surface. Our study sheds light on the mechanisms of methylammonium lead iodide perovskite formation and provides a pathway towards accessing the desirable morphology for the next generation of efficient perovskite solar cells. Experimental Section Materials: All reagents from commercial sources were used without further purification. PbI2 and PbCl2 with 99.999% purity, DMF, hydroiodic acid (HI) (57% in water), methylamine (CH3NH3) (33% in absolute ethanol), ethanol and diethyl ether were purchased from Sigma Aldrich.


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Synthesis of methylammonium iodide (CH3NH3I): 15 mL of hydroiodic acid and 14 mL of methylamine were stirred in the round-bottom flask at 0°C for 2 hours. A rotary evaporator was used to extract the white powder from the solution. In order to purify CH3NH3I, the white powder was re-dissolved in ethanol and re-precipitated in diethyl ether. A rotary evaporator was used to extract the purified white powder. The purification step was repeated three times. CH3NH3I crystals were collected and dried under a vacuum at 60°C for 48 hours. Thin-film fabrication: Glass substrates were sequentially cleaned by sonication in de-ionized water, acetone, and isopropyl alcohol, and dried in a nitrogen stream. A thin layer of PEDOT:PSS (Clevios™ PH 500, filtered with a 0.45 µm PVDF syringe filter – Pall Life Sciences) was first spin-coated on the pre-cleaned glass substrates and baked at 120°C for 40 minutes. PEDOT:PSS is typically used as a hole-transport layer in perovskite solar cells.9 While we have not reported device data in this paper, we wanted to mimic the processing conditions for active-layer processing in devices. CH3NH3I and PbX2 were dissolved in DMF at the desired molar ratios and stirred for 2 hours in the glovebox. Subsequently, the solution was spin-coated on top of the PEDOT:PSS-coated glass substrates at 4,000 rpm for 45 seconds and films were annealed at the pre-specified temperature for different times in the glovebox. After thermal annealing, the films were quenched to room temperature before they were placed in a custommade nitrogen-filled box for x-ray diffraction measurements. X-ray diffraction measurements: X-ray diffraction measurements were conducted on a Bruker D8 Discover diffractometer using Cu Kα radiation source (λ = 1.54 Å). The step size was 0.04° for the 2θ interval of 5° - 45°, while a 0.005° size step was used for the 2θ interval of 27.3° 29.3°.


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Scanning electron microscopy measurements: A low-voltage scanning electron microscope (FEI Verios 460 XHR) was used to image the morphology of fully-converted perovskite thin films. The accelerating voltage was kept at 5 keV to prevent beam damage to the specimens. ASSOCIATED CONTENT Supporting Information. X-ray diffraction traces, SEM images and optical micrographs of fully-converted perovskite thin films and integrated intensities of the (220) reflection of CH3NH3PbI3 as a function of annealing time for thin films having PbI2:CH3NH3I = 1:1 and PbCl2:CH3NH3I = 1:3 at the specified temperatures. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author [email protected] ACKNOWLEDGMENT The authors acknowledge the financial support from the National Science Foundation through grants ECCS-1549619 and CMMI-1537011 and partial support from the Andlinger Center for Energy and the Environment at Princeton University. REFERENCES (1)

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Crystalline Intermediates and Their Transformation Kinetics during the

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