Complete Suppression of Detrimental Polymorph Transitions in All

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Complete Suppression of Detrimental Polymorph Transitions in All-Inorganic Perovskites via Nanoconfinement Xiaoqing Kong, Kamran Shayan, Sophia Hua, Stefan Strauf, and Stephanie S. Lee ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.9b00322 • Publication Date (Web): 19 Mar 2019 Downloaded from http://pubs.acs.org on March 20, 2019

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ACS Applied Energy Materials

Complete Suppression of Detrimental Polymorph Transitions in All-Inorganic Perovskites via Nanoconfinement Xiaoqing Kong,† Kamran Shayan,§ Sophia Hua,† Stefan Strauf§ and Stephanie S. Lee*,† †Department §Department

of Chemical Engineering and Materials Science, Stevens Institute of Technology, Hoboken, NJ 07030, USA of Physics, Stevens Institute of Technology, Hoboken, NJ 07030, USA

KEYWORDS: All-inorganic perovskites, nanoconfinement, polymorph phase transition, stability, photophysics

ABSTRACT: Reducing the size of metal-halide perovskite crystals to the nanoscale has been demonstrated to stabilize highperformance metastable polymorphs at room temperature. Cesium lead iodide (CsPbI3), for example, typically exists as the insulating δ-phase at room temperature but can adopt the narrowband gap γ-phase when the crystal size is reduced to the nanometer length scale. Here we advance a fundamental understanding of the role of nanoconfinement in CsPbI3 polymorph stabilization through a combination of x-ray diffraction and temperature-dependent photoluminescence. Using a wet annealing method to directly form γCsPbI3 from solution in the cylindrical nanopores of anodized aluminum oxide, we discovered that nanoconfinement lowers the δ-γ solid-state phase transition temperature from 448 K in the bulk to 370 K. Once formed, nanoconfined γ-CsPbI3 crystals were found to be stable across the temperature range of 4 – 610 K and upon an unprecedented one year of air exposure at room temperature. Taking advantage of the nanoconfinement-induced suppression of phase transitions, we report for the first time a detailed analysis of electron-phonon interactions in γ-CsPbI3 via temperature-dependent photoluminescent measurements. In-depth analysis of the temperature-dependent peak broadening revealed electron-phonon interactions to be dominated by Fröhlich scattering, similar to that observed in inorganic-organic hybrid perovskite systems. Photoluminescence mapping further confirmed that nanoconfined γ-CsPbI3 crystals exhibit spatial uniformity on the tens of microns length scale, suggesting nanoconfinement as a promising strategy to form stable, high-performance perovskite films from solution for light-emitting and light-harvesting applications.

Introduction As power conversion efficiencies of solution-processed perovskite-based solar cells continue to increase beyond 20%,1,2 the development of strategies for stabilizing these devices against environmental degradation and detrimental polymorph transitions is critical to their commercialization. Inorganicorganic perovskites, such as methylammonium lead iodide (MAPbI3) and formamidinium lead iodide (FAPbI3), are highly hygroscopic and suffer from chemical decomposition at elevated temperatures. Replacing the volatile organic cation with a more stable inorganic ion, such as cesium, is an emerging strategy to improve the performance of light-harvesting materials in solar cells. Such inorganic perovskites exhibit excellent stability against moisture- and temperature-induced degradation,3,4 near-unity photoluminescence (PL) quantum yields, narrow emission peak widths,5 and negligible electron/hole trapping.6 All inorganic perovskite cesium lead iodide (CsPbI3) is currently being explored as the active layer of solar cells,7 red light-emitting devices8 and as top cells for perovskite/silicon tandem solar cells.9 However, the high-performance active phases of CsPbI3, namely the α-, β-, and γ-phases, are only stable at elevated temperatures.10 At room temperature, CsPbI3

exists in the inactive, non-perovskite -phase with a wide bandgap of 2.82 eV.11 Significant research efforts are thus underway to stabilize active CsPbI3 at room temperature through solution-based strategies, including compositional engineering methods, such as chemical doping with bromine or bismuth ions,9,12,13 dimensional methods to transform 3D CsPbI3 into stable 2D perovskites by introducing bulky ammonium cations,14 and polymer-induced surface passivation engineering.15 It is important to note that while the majority of these studies assume the active phase to be the cubic α-phase, more recent work suggests that in fact active CsPbI3 exists the orthorhombic γ-phase at room temperature.10,16,17 One of the most promising strategies to stabilize active CsPbI3 at room temperature is to reduce the size of CsPbI3 crystals. CsPbI3 or alloyed CsSn1-xPbxI3 quantum dots with diameters between 3 and 16 nm, for example, exist an active phase at room temperature.18–21 A variety of methods, such as the incorporation of solvents,22 small molecular ligands,23 and polymer additives,8 have been reported to stabilize an active phase of CsPbI3 via crystal size reduction. For example, the incorporation of hydroiodic acid (HI) into CsPbI3 precursor solutions was found to reduce the crystal size of CsPbI3 from the micron scale to ~100 nm in spun cast films.22 Similarly, a small amount (1.5 wt%) of sulfobetaine zwitterion added to 1

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Dry at RT for 10 min

Cool to RT

Dry Annealing: Anneal at 610 K

Spin coat

PbI2-CsI solution on AAO

Wet Annealing:

-CsPbI3

Cool to RT Anneal at 610 K

γ-CsPbI3

Scheme 1. Schematic of sample preparation in which co-solutions of PbI2 and CsI in DMF were deposited in AAO templates via spin coating and subsequently thermally annealed at 610 K in a nitrogen environment following two different methods. For the “dry annealing” method (top), samples were allowed to dry at RT for 10 min prior to transferring to the hot plate. In the “wet annealing” method (bottom), samples were immediately transferred from the spin coater to the hot plate with some solvent still present. Optical micrographs of the infiltrated AAO templates (diameter = 13 mm) after thermal annealing are provided.

CsPbI3 precursor solutions facilitated the formation of CsPbI3 crystals with average diameters of ~30 nm. The zwitterion ligands were found to impede the crystallization of CsPbI3 perovskite films via electrostatic interaction with the ions and during solution deposition.23 The addition of poly(ethylene oxide) to a CsPbI3 precursor solution can also reduce average CsPbI3 crystal size to ~100 nm in spun cast films, as the polymer chains impede the diffusion of the precursor ions during the crystal growth.8 In all of the above examples, X-ray diffraction analysis revealed these nanocrystals to adopt an active phase of CsPbI3, likely to be the γ-phase, at room temperature when stored in nitrogen. Here we explore the size-dependent thermodynamics of nanoconfined CsPbI3 across the temperature range of 4 – 610 K to advance a fundamental understanding of the role of nanoconfinement in stabilizing these systems against temperature-induced polymorph transitions. Nanoconfinement is a powerful strategy to select for metastable polymorphs of compounds, shift melting temperatures, and promote specific crystal orientations by reducing the average crystal size.24 We previously discovered nanoconfined MAPbI3 to be stable against humidity-induced degradation for over two years of storage in air.25,26 Nanoconfined MAPbI3, which can exist in at least three different phases depending on the temperature, also exhibits shifted polymorph transition temperatures compared to unconfined MAPbI3. Following these findings, we herein report on the use of nanoconfinement to stabilize the active perovskite phase of undoped CsPbI3 to temperatures as low as 4 K and up to one year of air exposure by shifting the Gibbs free energies of the solid-state phases and increasing the barrier to solid-state phase transitions. This unprecedented stability was exploited to characterize phonon interactions in these systems via temperature-dependent photoluminescence measurements.

Experimental Methods Perovskite CsPbI3 precursor solution synthesis. An equimolar 0.4 M mixture of PbI2 (Alfa Aesar, 99.999% purity) and CsI (Sigma-Aldrich, 99.999% purity) in anhydrous DMF (Sigma-Aldrich) was prepared in a nitrogen environment. The precursor solution was stirred overnight at 70 oC. Preparation of CsPbI3 films on substrates with varying extents of confinement. 50 l of the precursor solution was drop cast onto flat SiO2/Si substrates and nanoporous AAO

templates (Whatman®) with nominal pore diameters of 250 nm, respectively. The samples were subsequently spun cast at 2000 rpm for 30 s. In the dry annealing method, the as-spun samples were dried at room temperature for 10 min and then transferred to a hot plate set at various temperatures in the range of 350 – 610 K for 10 min. In the wet annealing method, the as-spun samples were immediately placed on a hot plate set at various temperatures in the range of 350 – 610 K for 10 min. In both methods, the samples were then cooled back to RT for further analysis. 2D x-ray diffraction measurements. 2D x-ray diffraction patterns (XRD) were collected on the samples using a Bruker AXS D8 DISCOVER GADDS diffractometer with VANTEC 2000 detector. The diffractometer was operated in reflection mode at 40 kV and 40 mA. Data was collected using an incident x-ray wavelength of λ = 1.5405 Å and at an incident angle of 3o. XRD patterns were collected in air at RT for 120 s. Samples were initially stored in an N2-filled container to prevent air exposure prior to data collection, and subsequently stored in ambient air for stability studies. Photoluminescence spectroscopy. Samples were placed in a closed-cycle cryogen-free cryostat with the accessible temperature range 4 – 300 K and ultralow vibration (attodry1100 from attocube). A green laser diode, emitting at 405 nm in continuous wave mode at a power of 7 W, was used for excitation. A diffraction-limited laser spot size of ~ 0.85 microns was achieved using a cryogenic microscope objective with numerical aperture of 0.82. The relative position between sample and laser spot was adjusted with a piezo-electric xyzstepper while 2D scan images were recorded with a 2D-piezo scanner (attocube). The PL emission from the sample was collected in a single mode fiber and dispersed using a 0.75 m focal length spectrometer equipped with a liquid nitrogen cooled silicon CCD camera. To realize hyperspectral images the signal was sent through optical band pass filters, as further detailed in Ref. 27.

Results and Discussion In a typical experiment, as illustrated in Scheme 1, a precursor solution containing stoichiometric CsI and PbI2 dissolved in N,N-dimethylformamide (DMF) was spin coated onto an anodized aluminum oxide (AAO) template comprising uniaxially aligned pores with characteristic diameters of 250 2


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nm. Using this method, approximately 50% of the pores were filled with no evidence of a capping layer (Figure S1). The samples were then either 1) dried at room temperature for 10 min before transferring to a hot plate at 610 K (i.e. the “dry annealing” method), or 2) immediately transferred from the spin coater to the hot plate with some residual solvent present (i.e. the “wet annealing” method). The samples were then annealed for 10 min and cooled to room temperature for further analysis. As displayed in the optical micrographs in Scheme 1, the final color of the infiltrated templates was different depending on the processing method. Samples processed via the dry annealing method exhibited a yellow color. X-ray diffraction (XRD) patterns collected on these samples (Figure 1A) revealed this color to correspond to the -phase of CsPbI3, consistent with previous literature reports that -CsPbI3 is the thermodynamically favored phase at room temperature.10 Interestingly, the samples processed via dry annealing exhibited a yellow color throughout the process, even when annealed at 610 K, 30 K above the bulk polymorph transition temperature to cubic α-CsPbI3. This observation suggests that nanoconfinement suppresses polymorph transitions in CsPbI3.

Figure 1. A) 2D XRD patterns collected at room temperature on a commercial AAO template infiltrated with CsPbI3 crystals annealed at 298 K and 610 K via the wet annealing method, respectively. B) 1D line traces along qxy = 0 Å-1 extracted from 2D XRD patterns collected on a commercial AAO template infiltrated with CsPbI3 crystals annealed at various temperatures via the wet annealing method. The XRD patterns were collected at room temperature approximately 1 hour after thermal annealing. Simulated powder patterns of γ-CsPbI3 and -CsPbI3 are provided at the bottom in black and red, respectively.16

In contrast, samples processed via wet annealing turned black within one minute on the hotplate and remained black upon cooling to room temperature. In this case, XRD patterns revealed this color to correspond to γ-CsPbI3 (Figure 1A). Scherrer analysis on this pattern confirmed the average crystallite size to be ~220 nm, approximately equal to the

characteristic pore diameter of the AAO scaffold of 250 nm (refer to Supporting Information). Surprisingly, we observed γCsPbI3 crystals confined in AAO templates to persist at room temperature for over one year of storage in air (with humidity values ranging from 30 to 95%), the longest yet reported shelflife of this compound (Figure S2). For comparison, in unconfined CsPbI3 crystals deposited on flat SiO2/Si substrates, the metastable γ-phase transitioned back to the -phase upon cooling within a period of two hours (Figure S3). Collectively, we observed that nanoconfined CsPbI3 at room temperature adopts the non-perovskite -phase when processed via dry annealing and the orthorhombic γ-phase when processed via wet annealing at 610 K. The phase that formed initially during drying was found to persist throughout processing and extended storage in air. Nanoconfinement thus suppresses polymorph transitions in CsPbI3, both from the - to γ-phase upon heating and from the γ- to -phase upon cooling. To determine if the suppression of polymorph transitions upon nanoconfinement is purely a kinetic phenomenon or also due to a shift of the relative Gibbs free energies of the polymorphs, we next performed a series of temperature-dependent experiments in which the annealing temperature, TH, in the wet annealing method was varied between 298 – 610 K. Figure 1B displays 1D line traces along qxy = 0 Å-1 extracted from 2D XRD patterns collected on CsPbI3-infiltrated AAO templates approximately 1 hr after wet annealing at TH. For reference, simulated powder patterns of both γ- and -CsPbI3 are provided.28 When wetannealed at TH = 350 K or below, nanoconfined CsPbI3 adopted the -phase. When wet-annealed at TH = 370 K or above, on the other hand, nanoconfined CsPbI3 adopted the γ-phase. In comparison, nanoconfined CsPbI3 crystals formed via the dry annealing method adopted the -phase for all TH values (Figure S4). These trends were observed regardless of the ambient humidity at the time of experimentation, indicating that waterinduced transitions are suppressed in nanoconfined CsPbI3 crystals. Based on the observations that 1) polymorph transitions are completely suppressed in nanoconfined CsPbI3 crystals and 2) γ-CsPbI3 can form at 370 K when crystallized directly from solution in nanoconfined spaces, 78 K lower than the bulk transition temperature of 448 K,10 we speculate that both the relative Gibbs free energies, G, of the polymorphs and the energy barrier of the polymorph transition are affected by nanoconfinement. The thermodynamic stability of a phase is determined by G, which in turn is a sum of the Gibbs free energy of the crystalline surface and of the bulk volume. Decreasing the crystal size results in an increase in the surface area-tovolume ratio, thereby increasing the relative contribution of the surface free energy as the extent of confinement increases. Correspondingly, Figure 2A displays our proposed temperature dependence of the Gibbs free energy of the respective phases in both the bulk and nanoconfined systems. As observed from Figure 2A, nanoconfinement increases the total Gibbs free energy of both of the phases, but the temperature at which the Gibbs free energies of γ and -CsPbI3 are equal decreases. Shifting polymorph phase transition temperatures upon decreasing the crystal size has been observed in other systems, including other perovskites, such as Al2O3,29 MAPbI325,30 and PbTiO3.31 It has been empirically observed across a wide range of inorganic oxides that crystal lattices tend to transform into a more symmetric structure as the crystal size approaches the tens to hundreds of nanometers.32 In these systems, atomic bonds can 3



A

γconfined δconfined

B

γbulk

δbulk 370

448

Temperature (K)

Gibbs Free Energy

have directional character that results in anisotropic lattice stretching along specific crystallographic directions. As the particle size decreases, the nature of the bonds becomes increasingly ionic, transforming crystals to higher-symmetry structures. In line with these findings, -CsPbI3 displays a higher degree of symmetry, with lattice dimensions of a = 8.6 Å-1, b = 8.9 Å-1, and c = 12.5 Å-1, compared to -CsPbI3, with lattice dimensions of a = 10.4 Å-1, b = 4.8 Å-1, and c = 17.8 Å1.10

Gibbs Free Energy

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

confined (T =370 K)

Ea,confined

bulk (T = 448 K)

Ea,bulk

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on CsPbI3 solid-state thermodynamics, AAO templates were modified with trichloro(phenyl)silane (TPS) to increase the surface hydrophobicity33 prior to CsPbI3 infiltration. TPS-AAO exhibited a water contact angle of 88.9o, compared to 41.8o for untreated AAO (Figure S8). Despite the large difference in surface energies of the templates, the temperature-dependent behavior of nanoconfined CsPbI3 in TPS-AAO templates was the same as that in unmodified templates. These results suggest that specific crystal-surface interactions are secondary compared to the role of nanoconfinement in shifting the relative surface versus volume Gibbs free energies of the crystals.

γ-phase δ-phase Reaction Coordinate

Figure 2. A) Proposed dependence of the Gibbs free energy on temperature as a function of phase for bulk and nanoconfined CsPbI3 crystals. B) Proposed reaction coordinate diagram of bulk and nanoconfined CsPbI3 crystals at their respective solid-state transition temperatures.

We also observed that once a polymorph is formed in nanoconfined pores, whether -CsPbI3 or -CsPbI3, it remains stable regardless of the thermal history applied to the sample (Figure S5). As previously mentioned, -CsPbI3 can only form in AAO nanopores at 370 K or greater. Once formed, however, lowering the temperature below 370 K does not induce a transition to the -phase. Nanoconfinement thus suppresses the - polymorph transition, suggesting that the energy barrier to solid-state transformation under nanoconfinement is much larger than in the bulk. Indeed, the conversion between the nonperovskite ( and perovskite (α, β,  forms of CsPbI3 requires a large change in the unit cell volume,10 which is likely prohibited by the confining AAO walls in our system. On the other hand, in situ temperature-dependent x-ray diffraction experiments confirmed that nanoconfined γ-CsPbI3 reversibly transitions to the α-phase above 610 K (Figure S6). Because the γ-α transition requires only a slight rotation of bonds, we expect this transition to be less sensitive to nanoconfinement. Based on these observations, Figure 2B displays a proposed Gibbs free energy reaction coordinate diagram of bulk and nanoconfined CsPbI3 crystals at Ttransition = 448 K and 370 K, respectively. At the transition temperature, the Gibbs free energies of the two phases are equal. For bulk CsPbI3 crystals, the activation energy barrier is relatively small. -CsPbI3 thus readily transforms to -CsPbI3 below the transition temperature, in line with our observations of CsPbI3 crystals deposited on flat SiO2/Si substrates. These samples turned black upon annealing at 610 K but transitioned back to the -phase within two hours of cooling to room temperature (Figure S7). For CsPbI3 nanocrystals, however, a prohibitively large energy barrier between the two polymorphs prevents solid-state phase transitions, kinetically trapping the phases over a broad range of temperatures. It is expected that crystal-surface interactions at the confining walls of the AAO templates influence the surface energy of the crystals. To examine the relative significance of this interaction

Figure 3. A) Normalized steady-state PL spectra of nanoconfined -CsPbI3 deposited in a commercial AAO template measured upon heating from 4 – 300 K. The arrows denote shifts in the emission peaks. B) Peak positions and C) FWHM values versus temperature extracted from Gaussian fits of the peaks displayed in (A). The solid black line in (C) is the fit of the FWHM of the PL peak to Eq. 1, which accounts for contributions from temperature-independent inhomogeneous broadening and Fröhlich coupling with LO phonons. Exploiting of the remarkable stability of -CsPbI3 in AAO, we performed temperature-dependent confocal micro-PL spectroscopy (μPL) on -CsPbI3 crystals to advance the current understanding of photophysics in this class of perovskites. Compared to a previous temperature-dependent magnetotransmission study on inorganic perovskites in which researchers kinetically trapped the active CsPbI3 phase by rapidly quenching samples from 623 to 4 K,34 in our experiments, special handling of the samples was not required to maintain the -CsPbI3 across a wide range of temperatures and upon air exposure. In this work, we investigate the PL spectra of nanoconfined -CsPbI3 under nonresonant (405 nm) excitation. Figure 3A displays the temperature-dependent PL spectra of nanoconfined -CsPbI3 obtained under a laser excitation of 405 nm during heating from 4 K to 300 K. Intensities are normalized to clarify the spectral emission peak shifts. As displayed in Figure 3A, the PL spectrum of the CsPbI3 at room temperature exhibits a single emission peak at 712 nm (1.74 eV). This energy position corresponds well with the band gap of -CsPbI3 as reported in literature.35 A second emission peak is present at temperatures below 70 K. For 4


D

SiO2/Si

E

P2

P1 3 µm

3 µm

P3

3 µm

F 3 µm

G

3

P1

2 1

Intensity (counts/s)

3 µm

3 µm

C

720 nm intensity

B

570 nm intensity

A

AAO

570 nm intensity

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720 nm intensity

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3

P2

H

2 1 80

P3

I

60 40 20 0 480

580

680

780

Wavelength (nm)

Figure 4. A) Top-view SEM image of CsPbI3 deposited in a commercial AAO template exhibiting average pore diameters of 250 nm via the wet annealing method. B and C) Spatially resolved μPL maps of the sample displayed in (A) collected at 300 K at excitation wavelengths of λ = 580 nm and 700 nm, respectively. Data was collected at room temperature approximately 1 hr after annealing at 610 K. D – F) Corresponding data for CsPbI3 deposited on flat SiO2/Si. G) Steady-state PL spectra of CsPbI3 deposited on SiO2/Si at spots P1, P2 and P3 marked in (E). The PL peaks of γ-CsPbI3 and δ-CsPbI3 are marked in red and blue, respectively.

convenience, hereinafter, the high-energy broad peak is denoted as Peak 1, and the low-energy peak is denoted as Peak 2. Dual PL emission peaks are a common feature in the low-temperature phase of methylammonium lead halide perovskite films.25,36–39 Both emissions have also been previously observed in the PL spectra of CsPbBr3 nanocrystals and single crystals.40,41 Observation of the excitation power dependences of the two emissions in this system suggested that both are excitonic in nature. The high- and low-energy peaks were attributed to the emission of free excitons and radiative recombination of trapped excitons, respectively.40 To extract the peak position and obtain the full-width at halfmaximum (FWHM) of each peak, we fitted the emission peaks with a Gaussian lineshape function. Figure 3B displays the positions of the two emission peaks extracted from the fittings versus temperature. As observed in the figure, Peak 2 blueshifted with increasing temperature from 1.65 eV at 4 K to 1.70 eV at 70 K before vanishing at higher temperatures. Likewise, the position of Peak 1 blue-shifted from 1.67 eV at 4 K to 1.74 eV at 298 K. No abrupt shifts in peak position were observed during heating, indicating that CsPbI3 remained in the -phase even at 4 K. A previous theoretical study found that the bandgap of MAPbI3 increases with lattice expansion upon heating, resulting in a blue shift of the PL peaks with temperature.42 We thus attribute the blue shift of Peak 1 with increasing temperature to thermal expansion of the CsPbI3 crystal lattice. Figure 3C displays the FWHM, labeled as  of Peak 1 of the steady-state PL spectra. Following the analysis presented by Herz and coworkers on inorganic-organic hybrid perovskite systems,43 we fitted  using the following model:    0 + γacT +

γLO exp

(

ELO

)

𝑘𝐵T ― 1

(1)

In this model, which has been widely applied to study electronphonon coupling in inorganic semiconductors,44 0 represents temperature-independent inhomogeneous broadening due to scattering from disorder and imperfections45 and varies with the shape, size and composition of crystals.46 The second term represents broadening due to charge carrier-acoustic phonon

interactions, where ac represents the strength of the coupling and is assumed to be a constant.47 The third term accounts for broadening due to charge carrier-longitudinal optical (LO) phonon coupling, also known as Fröhlich scattering. In this term, LO and ELO represent the strength of this coupling and the energy of the phonon based on its frequency, respectively.39,43 Below 40 K,  is independent of temperature. The temperatureindependent scattering, 0, was extracted from this regime to be 17 meV. This value is significantly larger than the smallest linewidth measured by our system, namely 0.1 meV in singlewall carbon nanotubes,48 indicating that 0 is not limited by scattering from the pump laser. The value of 17 meV is lower than that reported for MAPbI3, at 26 meV, and FAPbI3, at 19 meV.39,49 These results suggest that high-quality CsPbI3 crystals with relatively few defects form within AAO nanopores during processing. Furthermore, it is expected that charge carrieracoustic phonon interactions dominate in this low temperature regime and increase line broadening linearly with temperature. Our observation that  is independent of T below 40 K indicates that ac ≈ 0 meV/K, consistent with previous reports on organic-inorganic perovskites.39,49 Considering the data across the entire temperature range, we extracted values of LO = 38 meV/K and ELO = 11.7 meV. Figure 3C displays the predicted dependence of  on temperature using these parameters, in close agreement with our experimental data. In comparison, LO was measured to be 40 meV/K for both MAPbI3 and FAPbI3 using the same method of nonresonant PL excitation.40 These values are significantly lower than those of highly polar inorganic semiconductors such as GaN and ZnO, with reported values for LO around 1050 meV/K and 963 meV/K, respectively.50,51 Likewise, the LO phonon energy, ELO, we obtained for -CsPbI3 is close to those reported for MAPbI3 and MAPbBr3, at 11.5 meV and 15.3 meV, respectively.39,52 These values are slightly lower than those obtained for hybrid inorganic-organic perovskites using far-IR spectra determination of the LO phonon frequency.55 The reason for this discrepancy remains unclear but is likely due to difference in sample temperatures during measurements. Overall, the LO 5



phonon energies reported for metal-halide perovskite systems are much lower than those typically measured for various inorganic semiconductors. For example, ELO of GaAs and CdS were reported to be 30.4 meV and 38 meV, respectively.53,54 The large difference in LO phonon energy between perovskites and inorganic semiconductors results in the lower mobility of perovskites (MAPbI3 at 197 cm2V-1s-1) compared with inorganic semiconductors (GaAs at 7000 cm2V-1s-1).55 In the context of previously reported values for inorganic-organic hybrid perovskites, our results suggest that 1) the nature of the cation, whether organic or inorganic, has little effect on electronphonon interactions; 2) nanoconfinement does not significantly alter the photophysics of -CsPbI3, and 3) nanoconfinement improves the crystal quality compared to bulk processing, as we also observed in MAPbI3.25 Finally, taking advantage of the different PL peaks exhibited by - and -CsPbI3, we performed hyperspectral μPL mapping to examine the uniformity of nanoconfined versus bulk crystals. Specifically, spatially-resolved μPL maps were collected at room temperature with an attocube piezo-scanner and a laser spot size of ~850 nm on nanoconfined CsPbI3 crystals in an AAO template and on a CsPbI3 film deposited on a flat SiO2/Si substrate. Figures 4A – C display an SEM image of an AAO template infiltrated with -CsPbI3 via the wet annealing method and corresponding PL maps filtered around 580 ± 10 nm and 710 ± 10 nm, corresponding to the signal from - and -CsPbI3, respectively.56 As displayed in the Figure 4C, the PL response associated with -CsPbI3 was not present. The PL response of -CsPbI3, on the other hand, was uniform across the PL map, indicating that the nanoconfined crystals adopted only the phase. These results are in agreement with x-ray diffraction analysis that only the -phase is present in these templates. Figures 4D – F display the SEM and PL images for a CsPbI3 film deposited on a flat SiO2/Si substrate. These PL maps reveal strong spatial inhomogeneity in the -phase and -phase emission, with both present at room temperature. It is important to note that the data was collected 1 hour after annealing the film above the solid-state transition temperature of T = 610 K to convert crystals to α-CsPbI3, which in turn converted to βthen -CsPbI3 upon cooling. As previously mentioned, the transition from -CsPbI3 back to -CsPbI3 occurs over a period of minutes to hours upon storage at room temperature. After 12 hours, the sample displayed PL signal from only the -phase. The steady-state PL spectra as a function of wavelength at different locations (labeled P1, P2, and P3 in Figure 4E) of the CsPbI3 film deposited on SiO2/Si are presented in Figure 4G. In the PL spectra collected at location P1, a single peak located at ~580 nm is observed, corresponding to -CsPbI3. At location P2, two PL peaks are observed, indicating the co-existence of both -CsPbI3 and -CsPbI3. At location P3, only the PL signal corresponding to -CsPbI3 was observed. Nanoconfinement thus improves the material quality and the spatial uniformity of stable -CsPbI3 deposited via solution processing.

Conclusions In conclusion, we have discovered that nanoconfinement shifts the thermodynamics of CsPbI3, both lowering the - polymorph transition temperature and increasing the energy barrier of this solid-state phase transition. This shift in the Gibbs free energies of the polymorphs is a result of increasing the relative contribution of the surface free energy compared to that

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of the volume upon crystal size reduction. We exploited the prohibitively large energy barrier to polymorph transitions that completely suppresses the - transition in nanoconfined CsPbI3 crystals to temperatures as low as 4 K in order to examine the photophysics of -CsPbI3. Temperature-dependent PL measurements revealed Fröhlich scattering to be the dominant electron-phonon interaction, similar to other high-performing inorganic-organic hybrid perovskite systems. Furthermore, relatively small inhomogeneous broadening and uniform spatially-resolved μPL maps indicate that the nanoconfined crystals are of high quality over large areas. These crystals were found to be stable at room temperature in air for at least one year, the longest-yet reported shelf-life for this system. Collectively, these findings suggest nanoconfinement to be a promising strategy for the solution-phase processing of stable, high-performance perovskite films for applications in light emission and light energy harvesting.

ASSOCIATED CONTENT Supporting Information. SEM images, Scherrer analysis of crystallite size, optical micrographs, 2D XRD patterns and corresponding 1D XRD line traces on unconfined and confined CsPbI3 samples via wet and dry annealing methods. In situ temperature-dependent XRD measurements on confined CsPbI3. Contact angle measurements of water droplets on AAO templates without and with TPS treatment. (PDF) This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]

Author Contributions The manuscript was written through contributions of all authors.

ACKNOWLEDGMENT S.S. and S.S.L. acknowledge financial support under NSF award ECCS-MRI-1531237. The authors are also grateful for noteworthy support from PSEG to advance energy innovation at Stevens. The authors acknowledge the assistance of Dr. Chunhua Hu at the Department of Chemistry of New York University with 2D XRD experiments, and C.H. acknowledges support by the National Science Foundation under Award Numbers CRIF/CHE-0840277 and by the NSF MRSEC Program under Award Number DMR0820341. Research used microscopy resources within the Laboratory for Multiscale Imaging at Stevens Institute of Technology and the authors thank Dr. Tsengming Chou for assistance.

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Gibbs Free Energy

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1 µm

nanoconfined bulk -phase

Nanoconfined -CsPbI3

-phase

Reaction Coordinate


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Scheme 1. Schematic of sample preparation in which co-solutions of PbI2 and CsI in DMF were deposited in AAO templates via spin coat-ing and subsequently thermally annealed at 610 K in a nitrogen environment following two different methods. For the “dry annealing” method (top), samples were allowed to dry at RT for 10 min prior to transferring to the hot plate. In the “wet annealing” method (bottom), samples were immediately transferred from the spin coater to the hot plate with some solvent still present. Optical micrographs of the infiltrated AAO templates (diameter = 13 mm) after thermal annealing are provided. 661x191mm (72 x 72 DPI)


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Figure 1. A) 2D XRD patterns collected at room temperature on a commercial AAO template infiltrated with CsPbI3 crystals an-nealed at 298 K and 610 K via the wet annealing method, respectively. B) 1D line traces along qxy = 0 Å-1 extracted from 2D XRD patterns collected on a commercial AAO template infiltrated with

CsPbI3 crystals annealed at various temperatures via the wet annealing method. The XRD patterns were collected at room temperature approximately 1 hour after thermal annealing. Simulated powder patterns of γ-CsPbI3 and δ-CsPbI3 are provided at the bottom in black and red, respectively.16 287x337mm (72 x 72 DPI)



Figure 2. A) Proposed dependence of the Gibbs free energy on temperature as a function of phase for unconfined and nanoconfined crystals. B) Proposed reaction coordinate diagram of bulk and nanoconfined CsPbI3 crystals at their respective solid-state transition temperatures. 403x170mm (72 x 72 DPI)


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Figure 3. A) Normalized steady-state PL spectra of nanoconfined γ-CsPbI3 deposited in a commercial AAO template measured upon heating from 4 – 300 K. The arrows denote shifts in the emission peaks. B) Peak positions and C) FWHM values versus temperature extracted to Gaussian fits of the peaks displayed in A. The solid black line is the fit of the FWHM of the PL peak to Eq. 1, which accounts for contributions from temperature-independent inhomogeneous broadening and Fröhlich coupling with LO phonons. 390x352mm (72 x 72 DPI)



Figure 4. A) Top-view SEM image of CsPbI3 deposited in a commercial AAO template. B and C) Spatially resolved μPL maps of the sample displayed in (A) collected at 300 K at excitation wavelengths of λ = 580 nm and 700 nm, respectively. D – F) Corresponding data for CsPbI3 deposited on flat SiO2/Si. G) Steadystate PL spectra of CsPbI3 deposited on SiO2/Si at spots P1, P2 and P3 marked in (E). The PL peaks of γCsPbI3 and δ-CsPbI3 are marked in red and blue, respectively. 674x288mm (72 x 72 DPI)


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Table of Contents figure 618x313mm (72 x 72 DPI)


Complete Suppression of Detrimental Polymorph Transitions in All

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