Critical Intermediate Structure That Directs the Crystalline Texture and

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A Critical Intermediate Structure that Directs the Crystalline Texture and Surface Morphology of Organo-Lead Trihalide Perovskite Hao-Chung Chia, Hwo-Shuenn Sheu, Yu-Yun Hsiao, Shao-Sian Li, Yi-Kang Lan, Chung-Yao Lin, Je-Wei Chang, Yen-Chien Kuo, Chia-Hao Chen, ShihChang Weng, Chun-Jen Su, An-Chung Su, Chun-Wei Chen, and U-Ser Jeng ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b12378 • Publication Date (Web): 06 Oct 2017 Downloaded from http://pubs.acs.org on October 6, 2017

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ACS Applied Materials & Interfaces

A Critical Intermediate Structure that Directs the Crystalline Texture and Surface Morphology of Organo-Lead Trihalide Perovskite Hao-Chung Chia,1,# Hwo-Shuenn Sheu,2,# Yu-Yun Hsiao,3 Shao-Sian Li,3 Yi-Kang Lan,4 ChungYao Lin,1 Je-Wei Chang,1 Yen-Chien Kuo,2 Chia-Hao Chen,2 Shih-Chang Weng,2 Chun-Jen Su,2 An-Chung Su,1 Chun-Wei Chen3,* & U-Ser Jeng1,2,* 1

Department of Chemical Engineering, National Tsing Hua University, Hsinchu 30013, Taiwan

2

National Synchrotron Radiation Research Center, Hsinchu Science Park, Hsinchu 30076,

Taiwan 3

Department of Materials Science and Engineering, National Taiwan University, Taipei 10617,

Taiwan. 4

Materials and Electro-Optic Research Division, National Chung-Shan Institute of Science and

Technology, Taoyuan 32546, Taiwan #

these two authors contribute equally.

Corresponding Authors; *

E-mails: [email protected]

*

E-mail: [email protected]

ACS Paragon Plus Environment

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ABSTRACT

We have identified an often observed yet unresolved intermediate structure in a popular processing with dimethylformamide solutions of lead chloride and methylammonium iodide for perovskite solar cells. With sub-second time-resolved grazing-incidence X-ray scattering (GIXS) and X-ray photoemission spectroscopy, supplemental with ab initio calculation, the resolved intermediate structure (CH3NH3)2PbI2Cl2⋅CH3NH3I features in 2D perovskite bilayers of zigzagged lead-halide octahedra and sandwiched CH3NH3I layers. Such intermediate structure reveals a hidden correlation between the intermediate phase and the composition of the processing solution. Most importantly, the 2D perovskite lattice of the intermediate phase is largely crystallographically aligned with the [110] planes of the 3D perovskite cubic phase; consequently, with sublimation of Cl ions from the organo-lead octahedral terminal corners in prolonged annealing, the zigzagged octahedral layers of the intermediate phase can merge with the intercalated MAI layers for templated growth of perovskite crystals. Regulated by annealing temperature and the activation energies of the intermediate and perovskite, deduced from analysis of temperature-dependent structural kinetics, the intermediate phase is found to selectively mature first then melt along the layering direction for epitaxial conversion into perovskite crystals. The unveiled epitaxial conversion under growth kinetics controls might be general for solution-processed and intermediate-templated perovskite formation.

KEYWORDS: Organo-lead trihalide perovskite, intermediate-templated conversion, crystal growth kinetics, X-ray photoemission spectroscopy, grazing-incidence X-ray scattering


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1. Introduction Increasingly sophisticated device engineering has pushed impressively the power conversion efficiency (PCE) of organometal halide perovskite solar cells over 20% within a few years.1-4 Extensively reported recently is the critical importance of creating an appropriate intermediate phase to better control the subsequent formation process of perovskite crystalline thin films, hence improved surface morphology and crystal texture. Hiding behind the impressive achievement in PCE are the many proposed but unresolved intermediate structures and postulated conversion routes of the perovskite films for high-performance solar cells.5-18 With a center Pb cation that forms corner-sharing [PbX6] octahedra with halides X and a coordinated organic cation A, the APbX3 perovskite structure provides vast combinations for optoelectronic properties that match the requirements for solar cell applications, including the wide absorption range, low exciton binding energy, and tunable bandgap.1-4 With tailored film surface morphology and crystalline textures, heterojunction perovskite films can exhibit longer diffusion lengths and higher mobilities of the holes and electrons generated upon light illumination,1-4 hence improved charge transport to the device electrodes. Up to present time, methylammonium (MA) and formamidinium (FA) lead-halide perovskites with mixed I, Cl, and Br, are the most studied organometal halide perovskite solar cells.1-4 Sophisticated processing routes together with component replacements involving changes of the organic cation, halides, and metal atom, lead to greatly improved efficiency, stability, and/or environmental toxicity of organo-lead trihalide perovskite solar cells.5-7 The corresponding crystallization processes and mechanisms from the lead and organic salts, however, remains incompletely understood.8-11 Nevertheless, previous studies have established correlations between device performance and


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structures of the crystalline perovskite films, including surface morphology, crystal size, and crystallinity.1-4 For solution-processed perovskite films of MAPbX3, presence of crystalline intermediate phases was often evidenced with characteristic X-ray reflections of unidentified structures, prior to perovskite crystal formation. Intermediate phases are especially prominent in the spin-coated films from organic solutions containing mixed bromine/chloride salts,1-4,8-11 and could be further enhanced with solvent-engineering processing involved solvent-replacement during spincoating.12-13 Such intermediate phases were associated with enhanced crystal growth kinetics of perovskite observed using in situ grazing-incidence X-ray scattering (GIXS).8,14-18 Accumulated pieces of evidence15,19-23 suggest release of organic salts during the conversion of the intermediate phase to MAPbI3-xClx crystalline films.15,24-34 With unresolved intermediate structures, the proposed models and mechanisms for the intermediate-to-perovskite conversion are incomplete.1-4,24-34 In this study, combining temperature- and sub-second-resolved GIXS, X-ray diffraction (XRD), and X-ray photoemission spectroscopy (XPS), we unveil the structure and structural evolution of a widely observed yet unresolved intermediate phase in perovskite processing with typical dimethylformamide (DMF) solutions of lead chloride (PbCl2) and methylammonium iodide (MAI).1-4,8 We elucidate the crystal structure, composition, and fast formation kinetics of the Cl-incorporated intermediate phase over in situ annealing. The hence resolved intermediate structure MA2PbI2Cl2⋅MAI comprises 2D perovskite bilayers of zigzagged lead-halide octahedra and sandwiched CH3NH3I layers. Such structure manifests a long overlooked correlation between the intermediate phase and the processing solution.3,8 The structure and structural kinetics observed in this study elucidate an underlying epitaxial conversion route of the


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intermediated-templated formation of oriented organo-lead trihalide perovskite; the epitaxial conversion process under kinetics controls also provides hints on advanced processing for novel 2D perovskite materials.35

2. Results and Discussion 2.1 Structure of the Cl-incorporated intermediate phase. Shown in Figure 1a is the highly oriented GIXS pattern obtained (with the instrument setup shown in Figure S-1, Supporting Information, SI) for an especially thermally incubated pure intermediate phase L1 in the film spin-cast from a DMF solution containing MAI and PbCl2 (in 3:1 molar ratio). A corresponding GIXS pattern corrected for the rigorously defined scattering vector components qz and qr, respectively along the out-of-plane and in-plane directions of the film substrate orientation36 is given as Figure S-2a in SI. Inspired by the general crystal structure of 2D perovskites37-38 of A2BX4, reflections in Figure 1a are indexed with a monoclinic unit cell of MA2PbI2Cl2MAI, with the lattice parameters a = 6.30 Å, b = 29.35 Å, c = 8.49 Å, α = 90.0o, β = 92.1o, and γ = 90.0o; the structure is refined from the monoclinic P21/c structure of A2BX4 by incorporating MAI layers (Figure 1c,d). The indexed (0l0) and (021) reflections reveal polycrystalline texture with the b-axis lying preferentially along the film normal. A corresponding high-resolution X-ray diffraction (XRD) profile (Figure 1b), characterized by prominent (0l0) reflections of even l, supports a systematic absence originating from the 21-helical symmetry of the octahedra layers along the b-axis. Since the crystal orientation of the thin film (or crystal texture) could not be exhaustedly measured and taken into account in model simulation, the crystal structural refinement could hardly be done with minimization of the residual intensity differences between


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the simulated and experimental peaks.39 Alternatively, systematic structural adjustments convoluted with effects of preferred orientation and Debye-Waller factor were done with the model structure of the proposed unit cell (cf. Figure S-2), to best account for the observed peaks in the XRD and GIXS profiles (cf. the simulated profile in Figure 1b).14 We note that the model can be improved with the iodine ions of the MAI layer in the structural model (Figure 1c,d) deviating from the centrosymmetric position by ca. 0.3 Å with 50% probability to each side (Figure S-2d, SI). Such small structural deviation is, however, neglected in the following discussion without loss of generality. Overall, the structural model for the L1 phase features in a long crystallographic b lattice and MAI layers intercalating in-between the zigzagged lead-halide octahedra layers. A large crystal correlation length ξ (i.e. crystal grain size) of L1 ca. 250 nm along the film normal is estimated from the width of (020) reflection of the XRD profile in Figure 1b, using the Scherrer equation; the estimated crystal size is comparable to the spin-cast perovskite film thickness (ca. 350 nm) revealed from a cross-section TEM image of a similarly processed film (Figure S-3, SI). As illustrated in Figure 1c,d, the L1 phase comprises a bilayer of zigzagged octahedra of [PbI4Cl2]4− coordination, with iodine occupying all the sharing corners and chlorine the terminals corners. The MA positons in the proposed L1 crystal structure are further optimized by using the ab initio calculation with energy minimization40,41 on the basis of the lattice parameters and atomic sites of the octahedra determined from the X-ray data analysis. The energy-minimized crystal structure suggests that the L1 phase is stabilized by rich hydrogen bonding between the zigzagged octahedral layers and the surrounding ammonium groups of MA (Figure S-4, SI). The sandwiched MAI layers in the L1 phase play a critical role in interconnecting the 2D-strucutred octahedra into 3D perovskite crystals during high-temperature annealing, as detailed below.


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108

#

107 #

(060)

106

(110)P

(040)/(021)

(b) 107

108

XRD GIXS Simulation

(020)

#

106

#

(120)/(041) (130)

#

* (231) (082) (251)

(062)/(210)

* (081)

3

10

(130)

(100)

104

105

(042)

105

(011)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Relative Intensity

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2

10

4

6

8

104 10

3

10

2

10 12 14 16 18 20 22 24 -1

q (nm )

Figure 1. (a) An oriented GIXS pattern and (b) the corresponding XRD, GIXS (averaged from the 2D GIXS pattern), and simulated X-ray powder diffraction profiles, for the intermediate L1 phase incubated in a film spin-coated from a DMF solution containing MAI and PbCl2 (in 3:1 molar ratio). The representative reflections circled in (a) are indexed with the crystalline structure illustrated in (c) and (d) (orthogonal views), with a monoclinic unit cell of lattice vectors a, b, and c. The stars mark the reflections contributed partially by the FTO substrate. Marked with the symbol # in the XRD profile in (b) are the reflections of perovskite crystals MAPbI3-xClx emerged over a long X-ray data collection procedure. Squares mark the two reflections from unidentified crystals.


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2.2 PbCl2-based octahedra. We further elucidate the chlorine, iodine, and lead coordination environments of the bilayer-structured octahedra in the L1 phase with XPS. Shown in Figure 2a is a prominent Cl 2p core level spectrum observed for the sample of pure L1 phase, indicating rich Cl ions of a substantially lower binding energy (by 1.0 eV) than that of the precursor PbCl2 reported in literature42. The result suggests full conversion of PbCl2 into the L1 octahedra. In practice, without complexing with the greatly out-numbered MAI, PbCl2 of a poor solubility would have precipitated from the DMF sample solution prior to spin-coating.10 Further, the binding energy measured for the Cl ions at the octahedral terminal corners of the [PbI4Cl2]4− layers in L1 (Figure 1c), is 0.7 eV lower than that for MAPbCl3 perovskite crystals (Figure 2a) of all Cl ions taking the shared corners of [PbCl6]4− octahedra.42,43 The relatively lower Cl binding energy is attributed to additional hydrogen bonding of the terminal Cl ions of the [PbI4Cl2]4− octahedra layers with the surrounding MA+ moiety in the L1 phase. It is crucial for Cl ions to take the terminal positions, as they can dissociate and escape in prolonged high temperature annealing without disrupting the integrity of the 2D octahedral layer structure, as detailed below. The final product of MAPbI3−xClx showed no measurable Cl trace (Figure 2a), hence x ≈ 0. Furthermore, the skewed Pb 4f core level spectra in Figure 2b reveal an asymmetric nature of the lead environment [PbI4Cl2]4− in the L1 phase, as compared to the symmetrical peaks observed for the symmetric [PbI6]4− coordination of MAPbI3 or [PbCl6]4− in MAPbCl3 (Figure 2b). The slightly lower Pb 4f binding energy of the L1 phase than that of MAPbI3 suggests weaker association of Pb-halide in the [PbI4Cl2]4− octahedra than that in [PbI6]4−, in average. Since all the iodine ions in both cases are at similar sharing corners of the corresponding


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[PbX6]4− octahedra, the weakened Pb-X association is attributed to the terminal Cl ions of the L1 octahedra, having stronger hydrogen bonding with the MA moiety. The hydrogen bonding could result in an electron back-donation effect to the lead of [PbI4Cl2]4− for the reduced binding energy observed. In addition, the iodine 3d (I 3d) core level spectrum measured for the L1 phase shows a lower binding energy compared to that of [PbI6]4− of the final product MAPbI3−xClx (Figure 2c). Since the L1 phase comprises two different iodine environments of [PbI4Cl2]4− and MAI, the I 3d peak profile measured is deconvoluted into two peaks centered at 618.66 ± 0.02 and 618.25 ± 0.02 eV, respectively. The later peak, corresponding to the iodine ions of MAI, is 0.75 ± 0.03 eV lower than that of MaPbI3-xClx (Figure 2c), which is consistent with the difference measured for MAI and MAPbI3-xClx previously.43 All these XPS core level spectra of Cl 2p, I 3d, and Pb 4f synergistically elucidate the features of the local and electronic structures of the L1 phase (Figure 1c). Interestingly, a previous study44 speculated that [PbI4Cl2]4− octahedra could form flat PbI4 layer structures with two terminal chloride ions strengthening the PbI4 planes. Our GIXS and XPS results, however, reveal zigzagged [PbI4Cl2]4− octahedral layers, which would be more stable than planar ones due presumably to more interfaces for hydrogen bonding with the surrounding MAI.


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8000

14000

MAPbCl3

(a)

L1

-0.65 eV

Cl 2p

12000

Relative Intensity

10000

Relative Intensity

MAPbI3-xClx

6000

4000

(b)

MAPbCl3

MAPbI3

Pb 4f

L1

MAPbI3-xClx

0.40 eV

10000 8000 6000 4000 2000

2000 203 202 201 200 199 198 197 196 195 194

146

30000

(c)

25000

I 3d

MAPbI3-xClx

144

142

140

138

136

134

Binding Energy (eV)

Binding Energy (eV)

Relative Intensity

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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0.75 eV

L1

20000

15000

10000

5000

0 632

630

628

626

624

622

620

618

616

Binding Energy (eV)

Figure 2. XPS core level spectra of (a) Cl 2p, (b) Pb 4f, and (c) I 3d, measured for the perovskite films indicated. MAPbCl3 is prepared from MACl and PbCl2, MAPbI3 from MAI and PbI2, and the L1 phase and fully annealed MAPbI3-xClx (with x ≈ 0) from MAI and PbCl2. In (c), the peak for the L1 phase is fitted (red dotted curve) using the two profiles peaked respectively at 618.66 and 618.25 eV (marked by the vertical arrows) for iodine ions of [PbI4Cl2]−4 and MAI.

2.3 Formation and conversion kinetics. The maturing, melting, and conversion process of the L1 phase into perovskite at 110 °C annealing was captured in details, using subsecond-resolved (500 ms) GIXS. Fast development of the L1 phase features the first 100 s of annealing, which is accompanied by much slower perovskite formation (Figure 3a). In the subsequent 100 s, L1


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saturates with intensity (Figure S-5b, SI), which is followed by the major L1-to-perovskite conversion during t = 290 to 600 s. In the conversion, the major (020) and (040) reflections at q = 4.25 and 8.50 nm−1 of the L1 phase decay in synchrony without peak shifting. Concomitantly observed is the fast growth of the characteristic (100) peak of the high-temperature cubic phase of perovskite at a markedly different positon of q ≈ 10 nm−1 (Figure 3b).45,46 These features suggest melting and rearrangement of the L1 phase into perovskite, rather than direct stacking of the intermediate structure into perovskite crystals. Structural formation kinetics of the L1 phase is further extracted from the time-resolved GIXS data using the Avrami-Erofeev analysis;16,47 in which, the extent of a phase volume development is described by

α(t) = 1 − exp{−[k(t−t0)]n}

(1)

where the Avrami exponent n and rate constant k can be determined respectively from the slope and the intercept of the Sharp–Hancock presentation47,48 ln{−ln[1 − α(t)]} = nln (t – t0) + n ln k, with a properly chosen value of the induction time t0. We use the integrated peak intensity of the major peak (040)/(021) as the representative α of the L1 phase. Shown in Figure 3c are the temperature-dependent developments of the L1 phase at 100, 110, and 120 °C annealing. All these data can be fitted (Figure 3c) with a common n = 1.45 ± 0.05 and increasing k values with temperature, as summarized in Table 1. As there is little growth in crystal size (cf. Figure S-6, SI), the fitted n value larger than unity corresponds to an accelerated and nucleation-dominated process.47-49 Assuming the Avrami rate constant k follows the Arrhenius relation k(T) ∝ exp[−Ea/(RT)] with temperature T and gas constant R,47-49 we have further extracted an activation energy Ea = 72 ± 11 kJ/mol for the formation of the L1 phase (Figure 3d).


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7200 5400 3600

(002)

Perovskite

1800 0

Perovskite

(b)

5200 3900

(004)/(021)

2600 1300 0

70 60

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) t (s

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100 oC 110 oC o 120 C

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Ea = 72 kJ/mol

-5.0

ln(k)

ln (-ln (1-α))

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-2

Slope n = 1.45

-5.5 -6.0

-4

-6.5 -6 1

2

3

4

-7.0 0.0025

ln (t-t0)

0.0026 −1

0.0027 −1

T (K )

Figure 3. GIXS profiles measured in the (a) early and (b) late stages of 110 °C annealing, for a sample thin film spin-cast from MAI:PbCl2 (in 3:1 molar ratio), with (002) and (004)/(021) reflections of L1 and the major perovskite peak indicated. (c) The α(t) profiles for the L1 phase deduced from the (004)/(021) reflection respectively at 100, 110, and 120 °C annealing, fitted (dotted lines) using a common Avrami exponent n = 1.45 as indicated. (d) The temperaturedependent k(T) values in the Arrhenius plot are best fitted (dotted line) with Ea = 72 ± 11 kJ/mol.

Table 1. Best-fit values of the Avrami exponent n, rate constant k, and induction time t0 of the L1 phase and the perovskite crystals of the sample thin films at 100, 110, and 120 °C annealing.


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L1 (040) t0 (s) n k(10−3 s−1) Perovskite cubic (100) t0 (s) n1 k1(10−3 s−1) n2 k2(10−3 s−1)

100 °C

110 °C

120 °C

Ea (kJ/mol)

6.5 1.45 ± 0.04 1.80 ± 0.23

6.5 1.45 ± 0.01 4.39 ± 0.17

4.4 1.45 ± 0.02 6.37 ± 0.39

72 ± 11

100 °C

110 °C

120 °C

150 0.55 ± 0.03 0.020 ± 0.002 2.44 ± 0.06 1.34 ± 0.02

17 0.55 ± 0.01 0.56 ± 0.02 2.51 ± 0.09 2.50 ± 0.02

14 0.55 ± 0.02 8.9 ± 0.53 2.45 ± 0.22 7.50 ± 0.47

Ea (kJ/mol) 337 ± 21 103 ± 18

Shown in Figure 4a are the corresponding temperature-dependent growth behaviors of perovskite, deduced from Figures 3a and 3b on the basis of the integrated peak intensity of the primary reflection at q ≈ 10 nm−1 of perovskite for α(t). Unlike the monotonic growth mechanism with n = 1.45 of the L1 phase, perovskite formation can be characterized by two distinct stages with two markedly different n values. From the kinetic data in the early stage annealing (t = 17 to 140 s) at 110 °C, a significantly lower n1 (= na + ngd) = 0.55 is extracted,47−49 which is n

contributed mainly by the growth of the crystal size (or correlation length) ξ ∝ t g (with a fitted value of ng = 0.4 shown in Figure 4c) along the specific (100) crystallization direction, hence, the corresponding Euclidean dimension48 d = 1; this leaves a small contribution to primary nucleation na (= 0.15). After entering into the late stage where t > 290 s, the perovskite growth kinetics accelerates drastically to n2 ≈ 2.5 (Figure 4c), which is concomitant with the fast decay of L1 (Figure S-6). Since there is little perovskite size growth observed in the late stage (Figure 4c), the L1-toperovskite crystal conversion proceeds largely with n2 ≈ na for accelerated nucleation.48 We note that the increasingly more perovskite crystals formed from the L1-to-perovskite conversion gradually wash out the size contribution from the finite number of large and randomly oriented


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crystals formed in the early stage. From the temperature-dependent rate constants, the activation energy deduced for the perovskite formation in the early stage is Ea1 = 337 ± 21 kJ/mol (Figure 4b), much higher than that of the L1 phase (Figure 3d). Presumably, the higher energy barrier corresponds to the difficulty in structuring 3D-ordered perovskite from dispersed lead-halide octahedra in the as-spun film, under the competition with the L1 phase (organizing 2D-ordered perovskite octahedra with a lower Ea value). In contrast, a significantly lower activation energy of Ea2 = 103 ± 18 kJ/mol can be similarly extracted for the perovskite formation in the late stage (Figure 4b), via the L1 phase conversion with pre-ordered 2D perovskite octahedra. With completely suppressed formation of randomly oriented perovskite crystals in the early stage of annealing for pure L1 phase formation (via a slower heating process), especially enhanced cubic (100) oriented perovskite crystals from the L1-templated conversion can be achieved as shown in Figure 4d. Correspondingly, large crystal dimension (or correlation length) along the film normal of about 265 nm is deduced from the peak width of the final perovskite crystals after the annealing process (Figure S-7c, SI).45,46 The hence observed perovskite crystal size and orientation is comparable to that observed for the matured L1 phase prior to the conversion (Figure 1b). The superior formation kinetics of the Cl-incorporated intermediate phase plays a critical role in efficiently suppressing direct formation of randomly oriented perovskite in the early stage of annealing. This would not be the case with the Cl-free intermediate phase formed with MAI and PbI2 precursors in a similar spin-cast film; which exhibited parallel and comparable developments of the two different perovskite formation routes (i.e. direct vs. intermediateconversion), resulting in less aligned crystal orientations (Figure S-8, SI). The resolved formation kinetics shown in Figure 3 and 4 can provide guidelines in preparing organo-lead


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trihalide perovskite films for solar cells. For instances, annealing below the critical temperature 118 °C (the intersection of the two kinetic growth lines in Figure 4b), L1 phase can have increasingly faster formation kinetics to suppress direct formation of randomly oriented perovskite crystals. On the other hand, annealing above 100 °C provides better perovskite formation kinetics for improved crystallization rate and crystallinity. In practice, annealing at 110 °C has been commonly adopted on the basis of empirical wisdom.3

4 3

(a)

Ea2 = 103 kJ/mol

(b)

Ea1 = 345 kJ/mol

-4

1 0

ln(k)

ln(-ln(1-α))

-2

o

120 C 110 oC 100 oC

2

-1 -2 -3

n2 = 2.5

-4 -5

n1 = 0.55 2

3

4

5

-6 o

118 C

-8 -10 -12 0.0025

-6 6

0.0026 −1

ln(t-t0) ξ Relative Intensity

0.4

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1.2

0.8 0.6 0.4

50 n1 = 0

.5 5

0.15

Perovskite (100)cubic

(d) 30 q

z

(100) (200)

20

0.10

(100)

10

0.05

100

t (s)

200

0.0 400 600

L1 - (040)

(nm−1)

(100)

0 0

qr (nm−1)

0.2

0 10

0.20

1.4

1.0 n = 2 2.5

150

t

0.0027 −1

T (K )

Relative Intensity

200 (c)

Correlation Length ξ (nm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.00 -40 -20 0

20 40 60 80 100 120 140

Azimuthal angle Ψ (o)


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Figure 4. (a) Growth behavior of the α(t) profiles of the perovskite extracted from the integrated primary reflection intensity at q ≈ 10 nm−1, measured respectively at 100, 110, and 120 °C annealing. The profiles are respectively fitted (dotted lines) with common Avrami exponents n1 = 0.55 and n2 ≈ 2.50, in the regions before and after the L1-to-perovskite transition (marked at t ~ 290 s for 110 °C annealing). (b) Temperature-dependent Avrami rate constants in the two distinct growth stages respectively fitted with the Arrhenius relation, using the distinct Ea values indicated. (c) The typical growth behaviors of the intensity and the correlation length (or crystal size) ξ (fitted with t0.4 by a dotted curve) at 110 °C annealing. (d) Azimuthal intensity profiles of the primary (040)/(021) peak of the pure L1 phase and cubic (100) reflection of the perovskite formed at 110 °C, extracted from the 2D GIXS patterns in the inset and Figure 1a, respectively. The arrows mark the major and minor orientation preferences of the perovskite crystals along the film normal (Ψ = 90°) and 45° from it.

2.4 Epitaxial conversion model On the basis of the above structural information, we propose a process for the intermediatetemplated crystal growth of the oriented organo-lead trihalide perovskite thin films. As shown in Figure 5a, the terminal Cl-ions of the L1 octahedral layers destabilize to form MACl for sublimation in prolonged annealing;17 this is followed by interconnecting the 2D perovskite layers of activated octahedra [MAPbI4]4− with the sandwiched MAI layers along the layering direction. The zigzagged octahedral layers, of a 2D lattice crystallographically largely aligned with the [110] planes of the perovskite cubic phase (a = 6.220 Å; cf. Figure S-7, SI), can undergo epitaxial conversion to oriented 3D perovskite crystals (as detailed in Figure 5a); such process is consistent with the nucleation-dominated formation kinetics observed (with na ~ 2.5 for the L1to-perovskite conversion). As illustrated in Figure 5a, the L1 phase with [021] planes aligned to the film substrate can template the formation of perovskite crystals with cubic [100] planes in the same orientation; the crystal orientation could pass down to [110]-oriented tetragonal perovskite crystals (a = b = 8.914 Å and c = 12.716 Å; cf. Figure S-7, SI) after the cubic-to-tetragonal phase transition.45,46 Similarly, the [0l0] planes of the L1 crystals oriented parallel to the film substrate can convert into the cubic-[110] oriented perovskite crystals as shown in Figure 5b,


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corresponding to the oriented (100) peak at Ψ ~ 45° in Figure 4d. Furthermore, the SEM images taken for the films of pure L1 phase and that after fully annealing for perovskite (Figure 6) exhibit similar granule and pore features of micrometer sizes. The imaging result reveals that the final perovskite film inherits largely the morphological features (including surface coverage) of the L1 phase. Both the X-ray and imaging results illustrate the advantages of optimizing the intermediate phase prior to the intermediate-to-perovskite conversion process, for improved crystal size/texture and surface morphology.12,13 Our model in Figure 5 supports the proposed exchange of the chlorine and iodine ions in a recent report for successful fabrication of MA1−xFAxPbI3 lead halide perovskite.32 The model also predicts that without formation and sublimation of MACl, chlorine ions would continue to occupy the terminal corners of the octahedral layers of the L1 phase (Figure 5), hence, blocking the 2D-to-3D epitaxial conversion channel. Such inference is consistent with a recent observation17 that formation of perovskite could be greatly suppressed in a sealed sample environment with reduced Cl-evaporation during annealing. Our model further suggests that finite epitaxial conversion of 2D ordered octahedral layers of organo lead-trihalide is viable, provided that part of the mediating MAI layers in the L1 phase can be replaced by spacer layers of chemically inert organic moieties; which was evidenced lately.35,51,52 We note that the XPS results in Figure 2a could not exclude a possibility of residual chlorine existing deep beneath the perovskite film surface or near the CH3NH3PbI3−xClx perovskite/TiO2 interface, as evidenced recently.53 Previous theoretical calculations also suggested strong charge interactions between Cl and TiO2, resulting in enriched Cl ions at the MAPbI3−xClx and TiO2 interface; such Cl-TiO2 interface was expected to facilitate orientation alignment of the perovskite crystals and improve the TiO2 conduction band energy.54 The highly


17


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oriented L1 phase observed here may be a consequence of the proposed Cl-TiO2 charge interactions, via the terminal Cl of the layered octahedral of the L1 phase and the planar TiO2 layer (cf. Figure 5). In an advanced application, such chlorine-TiO2 interface has been applied as an excellent electron-selective interface to efficiently mitigate interfacial charge recombination and improve interface binding in planar perovskite solar cells.55 These Cl-TiO2 interface effects are complementary to the performance of the largely Cl-depleted perovskite active layer sitting on top the interface.54,55


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Figure 5. (a) A representative route for L1-to-perovskite phase transformation, with the [021] planes of the L1 phase parallel to the substrate (Right). After formation and sublimation of MACl in prolong high temperature annealing (wiggler arrows), the activated [MAPbI4]−4 octahedral layers (with a dihedral angle δ = 88° for the zigzagged octahedral plans, as indicated) can interconnect through the iodine ions of the MAI layers along the layering direction (Middle), for cubic perovskite crystals with [100] planes oriented to the substrate normal; the orientation passes down to tetragonal [110] oriented perovskite after the cubic-to-tetragonal phase transition during cooling (Left). (b) An alternative transformation route with the [0l0] planes of the intermediate L1 phase aligned along the substrate (Right). With similar transformation mechanism (Middle), the L1 octahedral layers transform to cubic [110] oriented perovskite crystals during annealing, followed by a transformation to tetragonal [100] oriented perovskite during cooling (Left).

Figure 6. SEM images (10 µm by 10 µm) taken for (a) amorphous as-cast, (b) pure L1 phase (annealed for 90s at 110 °C), and (c) fully annealed perovskite films, spin-cast from the sample solution of MAI and PbCl2. The latter two images exhibit similar features of polycrystalline granules and pores (marked with polygons) of micrometer sizes.

3. Conclusion We have elucidated an intermediate-templated epitaxial conversion process for oriented MAPbI3−xClx perovskite. The proposed intermediate structure comprises oriented 2D perovskite bilayers of zigzagged [PbI4Cl2] octahedra and sandwiched CH3NH3I layers, which structure unveils a hidden correlation to the optimized MAI:PbCl2 (3:1) molar ratio of the processing


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solution. With sublimation of the catalytic chlorine ions from the terminal corners during annealing, the bilayer-structured octahedra can undergo epitaxial conversion along the layering direction to the cubic phase of the perovskite crystals. The unveiled crystal structure and film morphology of the intermediate phase elucidate the critical importance of an optimized intermediate structure on solution-processed and intermediate-templated perovskite films. The kinetics controlled, epitaxial L1-to-perovsktie conversion process provides guidance on selection of processing parameters for improved crystalline texture and surface morphology of the perovskite films.

4. Experimental Section Sample Preparation. The sample thin films were spin-coated, from a DMF solution containing 40 wt% of MAI and PCl2 (in 3:1 molar ratio), onto FTO glass substrates (14 × 14 mm2) precoated with a compact TiO2 layer (ca. 200 nm thickness). The hence prepared perovskite precursor thin films (ca. 350 nm thicknesses as revealed from the TEM image in Figure S-3, SI) were then sealed in a N2 environment for in situ annealing with GIXS measurements. To trap the intermediate phase for X-ray diffraction and XPS measurements, the as-cast sample thin films were heated slowly in a chamber filled with N2, and cooled down a few minutes after the sample reached 110 °C.

Grazing Incidence X-ray Scattering. The as-cast sample thin films were respectively annealed at 100, 110, or 120 °C in a vacuum-tight chamber with flowing N2 for in situ GIXS at the BL23A SWAXS instrument56 of the Taiwan light source of National Synchrotron Radiation Research


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Center, Hsinchu. A 15-keV X-ray beam of an incident angle of 2.0° was used in GIXS data collection, with 500 ms time-resolution and a sample-to-detector distance of 192 mm. Data were collected using a CMOS flat panel detector C10158DK-3957 (10 × 10 cm2), of a pixel size of 50

µm.53 With the sample plane defined as the x-y plane and the incident X-ray beam in the x-z plane, the scattering vector q = (qx, qy, qz) is defined by q x = 2πλ −1 (cos β cos φ − cos α ) ,

q y = 2πλ−1 ( cos β sin φ ) , and qz = 2πλ−1 (sin α + sin β ) , with α and β for the incident and exit angles and φ for the scattering angle away from the y-z plane; qr = (qx2 + qy2)1/2 (detailed in Figure S-1, SI).57 The scattering peak positions were calibrated rigorously using the diffraction peaks from powders of Si, Ag-behenate, and LaB6.

Supporting Information GIXS and XPS measurements and supplemental data, XRD data analysis, ab initial calculation for energy minimization, and TEM cross-section image.

Acknowledgement. We thank W.-R. Wu for the help with GIXS measurements and Dr. O. Shih for discussions.

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Critical Intermediate Structure That Directs the Crystalline Texture and

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