3D heterojunction for high

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Interfacial engineering at the 2D/3D heterojunction for high-performance perovskite solar cells Tianqi Niu, Jing Lu, Xuguang Jia, Zhuo Xu, Ming-Chun Tang, Dounya Barrit, Ningyi Yuan, Jianning Ding, Xu Zhang, Yuanyuan Fan, Tao Luo, Yanlan Zhang, Detlef-M. Smilgies, Zhike Liu, A. Amassian, Shengye Jin, Kui Zhao, and Shengzhong (Frank) Liu Nano Lett., Just Accepted Manuscript • Publication Date (Web): 03 Sep 2019 Downloaded from pubs.acs.org on September 3, 2019

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Nano Letters

Interfacial engineering at the 2D/3D heterojunction for high-performance perovskite solar cells Tianqi Niu,#,1 Jing Lu,#,1 Xuguang Jia,3 Zhuo Xu,1 Ming-Chun Tang,4 Dounya Barrit,4 Ningyi Yuan,3 Jianning Ding,3 Xu Zhang,1,2 Yuanyuan Fan,1 Tao Luo,1 Yalan Zhang,1 Detlef-M. Smilgies,5 Zhike Liu,1 Aram Amassian,4,6 Shengye Jin,2 Kui Zhao*,1 and Shengzhong (Frank) Liu1,2 1Key

Laboratory of Applied Surface and Colloid Chemistry, Ministry of Education; Shaanxi Key Laboratory for Advanced Energy Devices; Shaanxi Engineering Lab for Advanced Energy Technology, School of Materials Science and Engineering, Shaanxi Normal University, Xi’an 710119, China. Email: [email protected] 2Dalian

National Laboratory for Clean Energy; iChEM, Dalian Institute of Chemical Physics, Chinese Academy of Sciences, Dalian, 116023, China. 3School

of Materials Science and Engineering, Jiangsu Collaborative Innovation Center of Photovoltaic Science and Engineering, Jiangsu Province Cultivation Base for State Key Laboratory of Photovoltaic Science and Technology, Changzhou University, Changzhou 213164, Jiangsu China. 4King

Abdullah University of Science and Technology (KAUST), KAUST Solar Center (KSC) and Physical Science and Engineering Division (PSE), Thuwal 23955-6900, Saudi Arabia. 5Cornell

High Energy Synchrotron Source, Cornell University, Ithaca, NY 14850, USA.

6Department

of Materials Science and Engineering, North Carolina State University, Raleigh, NC, 27695, USA.

ACS Paragon Plus Environment

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ABSTRACT: Perovskite solar cells based on two-dimensional/three-dimensional (2D/3D) hierarchical structure have attracted significant attention in recent years due to their promising photovoltaic performance and stability. But to obtain a detailed understanding of interfacial mechanism at the 2D/3D heterojunction, e.g. the ligand-chemistry dependent nature of the 2D/3D heterojunction and its influence on charge collection and the final photovoltaic outcome, is not yet fully developed. Here we demonstrate that the underlying 3D phase templates growth of quantum wells (QWs) within 2D capping layer, which is further influenced by the fluorination of spacers and compositional engineering in terms of thickness distribution and orientation. Better QW alignment and faster dynamics of charge transfer at the 2D/3D heterojunction result in higher charge mobility and lower charge recombination loss, largely explaining the significant improvements in charge collection and open-circuit voltage (VOC) in complete solar cells. As a result, 2D/3D solar cells with power-conversion efficiency (PCE) of 21.15% were achieved, significantly higher than the 3D counterpart (19.02%). This work provides key missing information on how interfacial engineering influences the desirable electronic properties of the 2D/3D hierarchical films and device performance via ligand chemistry and compositional engineering in QW layer. KEYWORDS: 2D/3D heterojunction, perovskite solar cell, ligand chemistry, interfacial mechanism, high-performance


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Introduction Organic-inorganic halide perovskites have attracted tremendous research interest in the fields of photovoltaics, photonics and optoelectronics due to their excellent properties such as high absorption coefficient,1,2 long exciton diffusion length,3 and tunable direct bandgap.4 This class of three-dimensional (3D) perovskites has revolutionized the field of photovoltaics by achieving a certified power conversion efficiency (PCE) of up to 25.2%.5 Despite these advantages, perovskites are susceptible to the external environment such as moisture, heat and irradiation, the degradation and phase conversation of perovskite films therefore hinder the device performance. In contrast to the 3D counterparts, the two-dimensional (2D) Ruddlesden-Popper (RP) perovskites with a general formula of (RNH3)2(A)n-1MnX3n+1 have recently attracted wide researches due to their tunability of optoelectronic properties and, more importantly, ambient stability.6-10 The hydrophobic nature of the spacing layer as well as the dense packing crystal structure prevent perovskites from the direct contact of moisture and therefore the solar cells based on 2D perovskites could remain stable in ambient environment for thousands of hours. Building on the added benefits of high stability, these 2D perovskites were further incorporated onto the surface of 3D components, as a capping layer to promote the stability of the underneath 3D perovskite phase without significant scarification of device performance. The 2D/3D hierarchical structure was firstly constructed for 3D MAPbI3 (MA = methylammonium). For example, solar cells based on the (PEA)2(MA)4Pb5I16/MAPbI3 (PEA = phenylethylammonium) 2D/3D hierarchical structure exhibited a strong moisture hindrance (75% RH).11 Huang et al. demonstrated a (BA)2PbI4/MAPbI3 (BA = butylammonium) stacking structure with much enhanced thermal stability and a high PCE of 19.89%.12 Recently, the stabilizing benefits of 2D capping layer have been extended to 3D FA-alloyed (FA=formamidinium) systems.13


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Nazeeruddin et al. obtained a PCE of 20.75% for the PEA2PbI4/Cs0.1FA0.74MA0.13PbI2.48Br0.39 2D/3D hierarchical structure, which retained 85% of the initial value after illumination for 800 hours at 50 ºC.14 Wang et al. found larger Fermi-level splitting and reduce non-radiative recombination losses at the PEA2PbI4/Cs0.05(FA0.83MA0.17)0.95Pb(I0.83Br0.17)3 heterojunction, which help increase open-circuit voltage (VOC).15 Fluorinated spacer cations, which are known to be highly hydrophobic,16 have recently been applied in 2D/3D hierarchical structure with a PCE of

20.00%

and

excellent

stability.17

Beyond

that,

other

organic

ligands

like

cyclopropylammonium (CA+) and 5-aminovaleric acid (AVA+) with optimized concentration have also been introduced to the 2D/3D hierarchical structure with simultaneously enhanced stability.18,19 Progress has been made with the 2D/3D hierarchical structure based on different spacer cations to achieve high photovoltaic operational stability under ambient air conditions. For neat 2D perovskites, much has been achieved in determining the nature of films, including thickness distribution, orientation, and charge transport between quantum wells (QWs).20-22 But to obtain a detailed understanding of interfacial mechanism at the 2D/3D heterojunction, e.g. the ligandchemistry dependent nature of the 2D/3D heterojunction and its influence on charge collection and the final photovoltaic outcome, is not yet fully developed. The ligand chemistry is significant importance to the thickness distribution and orientation of 2D perovskites,23 which is expected to play a key role in charge transfer at the 2D/3D heterojunction and charge transport/extraction in a complete solar cell. Thus, interfacial engineering via tuning ligand chemistry of 2D perovskites and its impact on the 2D/3D heterojunction as well as the final photovoltaic outcome become an interesting yet under-explored direction. Such a study is critical for identifying and improving the built-in voltage and charge recombination losses, as well as variations of photovoltaic device


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figures of merit and ambient stability of solar cells, with further realization of high-performance and air-stable perovskite solar cells for widespread employment. Herein, we sought to increase the systematic understanding of how interfacial engineering via fluorinating spacer ligands or compositional tuning influences charge transport/extraction and photovoltaic performance of FAPbI3 based 2D/3D solar cells. The results indicate a templated growth of QWs with small n values by the underlying 3D FAPbI3 layer, which induces packing and orientation different from that within neat QW films or single crystals. The phase-pure (n = 2) QWs are observed for the aliphatic ligands of n-butylammonium (BA) and 4, 4, 4trifluorobutylammonium (FBA). While phase impurity of QWs (n = 1 and 2) are observed for the aromatic ligands of phenylethylammonium (PEA) and 4-fluorophenylethylammonium (FPEA), and can be improved via composition engineering of introduced ligands. Interestingly, fluorination of ligands was found to lead to more randomly orientated wells. The charge transfer from QWs to the underlying 3D FAPbI3 layer at the 2D/3D heterojunction was further confirmed, with faster kinetics for the PEA-based phase-pure (n=2) QWs. Specifically, the QWs having better orientation and faster charge transfer kinetics at the 2D/3D heterojunction are beneficial for carrier mobility and suppression of charge non-radiative recombination for 2D/3D solar cells. As a result, a low loss-in-potential and a high charge collection efficiency were obtained for the 2D/3D solar cells based on PEA-based phase-pure (n=2) QWs, where carrier transport and passivation achieve an optimal balance. The solar cells show a JSC of 24.20 mA cm-2, a FF of 76.6%, a VOC of 1.14 V and a PCE of 21.15%, among the highest values for planar 2D/3D perovskite devices, and much higher than the 3D counterpart (PCE = 19.02%). Furthermore, excellent ambient stability was also confirmed for the 2D/3D structure devices, demonstrating the indispensability of the 2D phase to the device stability.


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Results and discussion 2D/3D hierarchical structure. We have first realized the 2D/3D perovskite hierarchical structure as illustrated in Fig. 1a. The control FAPbI3 was prepared from a precursor solution of formamidinium iodide (FAI, HC(NH2)2I) and PbI2 with a molar ratio of 1:1 under the concentration of 1.2M (see details in the Experimental section). The FAPbI3+PbI2 films were deposited from FAI:PbI2 (1:1.08). Isopropyl alcohol (IPA) solutions containing iodized spacers were introduced onto the thermal annealed FAPbI3+PbI2 films. The slightly excess PbI2 ensured the complete reaction with the iodized spacers, which had negligible influence on the device performance (Fig S1). Fig. 1b illustrates the spacer ligands used: phenylethylammonium (PEA), n-butylammonium

(BA),

4,

4,

4-trifluoro

butylammonium

(FBA)

and

4-fluoro

phenylethylammonium (FPEA), as well as PEA:FA mixtures with different weight ratios. Further thermal annealing leads to the reaction of PbI2 with spacer cations (Fig. S2) and generates 2D/3D hierarchical structure (Fig. 1c). The 2D perovskites follow the formula (L)2(FA)n-1PbnI3n+1, where L is the spacers FBA, FPEA, BA or PEA. The morphologies of the 2D/3D perovskite films were evaluated using scanning electronic microscopy (SEM) and atomic force microscopy (AFM). A thin and compact capping layer was observed from the cross-sectional SEM images (Fig. 1d). The plan-view SEM images show “plate-like” morphology with feature sizes on the order of hundreds of nanometers for the 2D/3D films (Fig. 1e), which is in contrast to the 0.5-1.0 P

4 3 grain-like features of the control 3D

film. In greater detail, AFM topography images show thin stacks of 2D planes on the surface of the 2D/3D films (Fig. 1f). Interestingly, the root-mean-square roughness (RMS) shows a significantly decrease in the presence of 2D capping layer, suggesting better surface quality. We


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images of the control and PEA based 2D/3D films. (f) Atomic force microscopy (AFM) images of the control and PEA based 2D/3D films showing root-mean-square roughness (RMS). (g) Plan-view SEM images of the FBA, FPEA, BA and PEA:FA(1:1) based 2D/3D perovskite films. Texture and orientation of quantum wells. The ligand-chemistry dependent QW texture and orientation within 2D capping layer were assessed using grazing incidence wide-angle X-ray scattering (GIWAXS) analysis. The GIWAXS snapshot for the control 3D film exhibits scattering peaks at q = 9.9 nm-1 and q = 8.5 nm-1, corresponding to the perovskite 3C (100) phase and the hexagonal (non-perovskite) 2H (100) phase, respectively (Fig. S3).24 The unexpected appearance of the 2H (100) phase is due to phase instability during shipment to the synchrotron and measurement under ambient conditions, which is significantly suppressed in the 2D/3D films. The underlying 3D phase templates growth of QWs within 2D capping layer. For example, the 2D capping layer exhibits diffraction arc of (PEA)2PbI4 (2.6 and 3.7 nm-1) (Fig. 2a), which is highly contrast to the features in the neat (PEA)2PbI4 film and (PEA)2PbI4 single crystal. The neat (PEA)2PbI4 phase exhibits the repeating periodic signature of the (00h) (h =1, 2, 3…) planes in the out-of-plane direction and (11 ) in the in-plane direction, a signature of the parallel orientation of QWs (Fig. 2b).25 The line-profiles of the three snapshots highlight the diffraction difference between the (PEA)2PbI4 QWs in the neat film/crystal and 2D/3D films (Fig. 2c). The observation implies that QW in the 2D/3D films does not follow its intrinsic parallel packing, which is ascribed to templated growth by the underlying 3D lattice (Fig. 2d). This templated growth further decreases the barrier for charge transport along the z-axis within 2D capping layer, and is beneficial for improving short-circuit current (JSC) in 2D/3D solar cells.


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Nano Letters 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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intensity versus q for the three cases. (d) The growth orientation of quantum wells within the 2D/3D film, (PEA)2PbI4 film and (PEA)2PbI4 single crystal. (e) Ex situ GIWAXS patterns of the FBA, FPEA and BA based 2D/3D films. (f) The corresponding intensity versus q for the diffraction features of the control and 2D/3D films. (g) The pole figures for the four 2D/3D films showing the different intensity and broadness in Azimuth angles of the n = 2 phase. Ligand chemistry plays an important role in the thickness distribution and orientation of QWs within the 2D/3D films. The diffraction arcs of QWs are observed for all 2D/3D films (Fig. 2e) without periodic feature for the highly parallel orientation, suggesting a general property of 3Dinduced templated growth of QWs. The aliphatic spacers BA- and FBA-based 2D/3D films exhibit phase pure (n = 2) QWs (q = 3.1 nm-1),22 while aromatic spacers PEA- and FPEA-based cases show phase impurity (co-existence of n = 1 and n = 2) of QWs (Fig. 2f).26 The azimuth angles of the n = 2 phase for the four 2D/3D films were investigated to evaluate the orientation randomness of QWs. Higher intensity at 90º is correlated with the formation of crystals with more out-of-plane texture. We observed dramatic reduction in the presence of anisotropic rings for the PEA- and BA-based 2D/3D films in contrast to the fluorinated ones, implying higher randomness of QWs when fluorinating spacers. Charge dynamics at heterojunction. The 2D/3D heterojunction enables built-in band alignment within in thin films. The linear absorption spectra exhibit a bandgap (Eg) excitonic peak for 3D phase (ca. 1.54 eV) and a peak for n = 2 QWs (ca. 2.18 eV) for all films (Fig. 3a and Fig. S4).27 Additional peak for n = 1 QWs was observed within PEA- and FPEA-based films which can be vanished via compositional engineering. Photoluminescence (PL) spectra exhibit 2fold increase in emission intensity of the 3D phase within the 2D/3D films under front-side excitation (film surface) in contrast to back-side excitation (glass side), suggesting a suppression


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of trap formation at the 2D/3D heterojunction (Fig. 3b). This enhanced radiative efficiency is reported to correlate with the gain of electron volts in the quasi-Fermi level splitting.28 First-principles calculations were further performed to evaluate electronic properties at the 2D/3D heterojunction (Fig. 3c). The density of states was calculated for the representative (PEA)2PbI4, FAPbI3 and their interface. The determined band gap (Eg) for the bulk (PEA)2PbI4 and FAPbI3 are 2.08 and 1.47 eV, respectively, with higher value for the former due to larger charge localization.29 The 2D/3D heterojunction shows 0.05 eV higher in the Eg compared to the bulk FAPbI3. The valence band (VB), referenced to the vacuum level, shows -4.80 and -5.13 eV for the bulk (PEA)2PbI4 and bulk FAPbI3, respectively. Meanwhile, the conduction band (CB) decreases by 0.94 eV for the bulk FAPbI3 in contrast to the bulk (PEA)2PbI4 phase (-3.66 vs. 2.72 eV). These observations are also confirmed in the presence of spin-orbit-coupling (SOC), and indicate a built-in band alignment at the 2D/3D heterojunction, which is beneficial for the internal charge separation and charge transfer at the junction. The influence of ligand-chemistry on charge dynamics at the 2D/3D heterojunction was further evaluated. Ultrafast transient absorption (TA) spectroscopy was performed on thin films with a femtosecond laser pulse. The photo-induced changes in the absorption 5U 6 spectra were then probed with a time-delayed laser-generated white light probe pulse.20,30,31 Fig. 3d illustrates the TA spectra at various delay times for all 2D/3D films. Phase-pure (n = 2) and phase impurity (n = 1, 2) QWs were observed for the aliphatic spacer-based and aromatic spacer-based 2D/3D films, respectively, consistent with the GIWAXS observation. A blue shift of n = 2 QW was found when triple-fluorinating BA (2.17 vs. 2.21 eV) (Fig. 3e), due to polarization-assisted decreased disparity of dielectric constant between the inorganic and organic layer.32 By contrast, low polarization for the mono-fluorinated spacers combined with symmetrical configuration


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leads to negligible change of excitonic resonance.33 Fig. 3f shows the typical dynamical charge dynamics for the representative PEA-based 2D/3D films (see Fig. S5 for other films). The formation of 3D bleaching is accompanied with a fast decay of QW bleaching, which is previously ascribed to the electron transfer from QWs to 3D phase, accompanied by the hole transfer in the opposite direction.20,34,35 The charge transfer at the 2D/3D heterojunction is comparable in the PEA-, FPEA-, BA-, and FBA-based films, while faster in the FA:PEA (1:1)based film (Fig. 3g). By contrast, there was no difference in 3D bleaching under the back excitation (Fig. S6 and S7). We can draw a conclusion from above observations that ligand chemistry of spacers holds the key to the QW nature in terms of thickness distribution and orientation. Better out-of-plane orientation of QWs was achieved for the spacers without fluorination. Phase pure QWs tend to be formed for the aliphatic-spacer based film, which can be also improved via compositional engineering. Charge transfer spontaneously occurs at the 2D/3D heterojunction (Fig. 3f) within sub-ps timescale, indicating the extremely high efficiency of overall charge transfer process. The FA:PEA (1:1)-based film exhibits slightly faster charge transfer dynamics. However fluorination of spacers plays negligible influence on the charge transfer dynamics, which instead, possess impact on passivation at the 2D/3D heterojunction and charge transport within the 2D capping layer as discussed below.


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Nano Letters 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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the glass illuminated from the front and back sides. (c) DOS analysis of the (PEA)2PbI4, FAPbI3 and their interface. (d) Transient absorption (TA) spectra at different delay times of the control and different 2D/3D films under front-excitation. (e) TA spectra at t = 2 ps with front-excitation for the different films. (f) The dynamic evolution of the bleaching recovery for the PEA based 2D/3D films. (g) The TA kinetics of the 3D bleaching for series of films with front-excitation. (h) Schematic model of the charge transfer between the 2D and 3D phases. Ligand-chemistry dependent passivation. The ligand-chemistry dependent passivation at the 2D/3D heterojunction was further evaluated. The trap density and charge mobility were measured for all devices with FTO/c-TiO2/perovskite/PCBM/Ag architecture (Fig. S8). The trap density was determined using the following equation:36

n

where

0

=

2

is the vacuum permittivity,

0

(1)

2

e

is the relative dielectric constant, VTFL is the onset voltage

of the trap-filled limit region, e is the elementary charge, and L is the distance between the electrodes.35 The electron mobility was further extracted using the Mott–Gurney Law:36

=

8 9

0

3 2

(2)

where JD is the current density, and V is the applied voltage. The trap densities are 1.70 ± 0.28, 1.23 ± 0.26, 0.87 ± 0.21, 0.91 ± 0.24, 0.65 ± 0.20 and 0.41 ± 0.18× 1016 cm-3 for the control 3D and the FBA-, FPEA-, BA-, PEA- and PEA:FA(1:1)-based 2D/3D films (Fig. 4a), respectively. Lower trap densities were obtained in the films without fluorination, in particular for the PEA:FA(1:1)-based film, suggesting more efficient passivation at the heterojunction and less non-radiative recombination losses.12 The mobility decreases from ca. 4.4 to 1.9-2.5 cm2 V-1 s-1


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for the fluorinated 2D/3D films in contrast to the control 3D film. The reduction in electron mobility for the fluorinated cases is agreement with higher orientation randomness of QWs as observed from GIWAXS analysis. By contrast, we achieve comparable mobility in the BA-, PEA-based films with the control 3D film, and even higher mobility in the PEA:FA(1:1)-based 2D/3D case. The PEA:FA (1:1)-based film with higher charge mobilities and lower trap densities is expected to deliver more efficient charge collection and higher fill factor (FF) of solar cells. The charge recombination was investigated using time-resolved PL spectroscopy (TRPL) (Fig. 4b), which was performed for the films on glass substrate and the correlated parameters were fitted by a bi-exponential equation:36,38 =

( 1

0)/ 1

+

( 2

0)/ 2

(3)

where $1 and $2 are the slow and fast decay time constants, respectively, and A1 and A2 are their corresponding decay amplitudes. The average lifetime 5Yave) values decrease from ca. 172 to ~130-137 ns for the fluorinated 2D/3D films in contrast to the control 3D one (Table S2).33,39,40 The Yave values increase to ca. 237, ca. 314 and ca. 270 ns for the BA-, PEA- and PEA:FA(1:1)based 2D/3D films, respectively. The PL mapping exhibits higher uniformity for the 2D/3D films without fluorination, particularly of the PEA:FA(1:1)-based one (Fig. S9). The higher charge lifetime, a signature of lower recombination loss, is due to better passivation. In other words, the ligand chemistry and compositional engineering determine how QWs form at the 2D/3D heterojunction which in turn plays a critical role on the passivation at the heterojunction and charge transport within the 2D/3D films. We achieved lower charge recombination loss for the spacers without fluorination, and higher charge mobility for QWs with better out-of-plane


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orientation. The PEA:FA(1:1)-based 2D/3D film achieves a balance between passivation and charge transport. Photovoltaic performance. Planar devices were fabricated with the n-i-p architecture of glass/FTO/c-TiO2/perovskite/Spiro-OMeTAD/Au. Fig. 4c illustrates the PCEs of all devices. The corresponding device parameters are summarized in Table S1. The best control device shows PCEmax of 19.02% (average PCE of 18.10 ± 0.46% for 40 cells), with a VOC of 1.06 V, a JSC of 23.99 mA cm-2 and a FF of 74.8%. The fluorinated 2D/3D perovskites show a PCE drop in contrast to the control one, due to lower charge mobilities and higher charge recombination losses. By contrast, high PCEmax values of 19.50%, 19.84%, and 21.15% were obtained for the BA-, PEA- and PEA:FA(1:1)-based 2D/3D devices, mainly due to a significant enhancement in the VOC (1.14 vs. 1.05 V) and a slight increase in the FF (76.6% vs. 74.5%) in contrast to the control (Fig. 4d and Table S1). This performance is also among the highest for the 2D/3D hierarchical perovskite devices.41-46 The integrated current density from the external quantum efficiency (EQE) matches the JSC values measured under AM 1.5G one sun illumination (Fig. S10). Hysteresis is also significantly improved within PEA:FA (1:1)-based 2D/3D device partially due to suppressed trap formation,47-48 with hysteresis indexes significantly decreasing from 8.9% to 2.1% in contrast to the control case (Fig. 4e and Table S3). The stabilized power output analysis under continuous AM 1.5-G, 1-sun illumination at a fixed maximum power point (MPP) voltage for 300 s shows much faster rise of output efficiency and better illumination stability for the PEA:FA (1:1)-based 2D/3D device in contrast to the control one (Fig. 4f and Fig. S11).8


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Mobility Trap density

6

FBA

1

3

10

2

0 Control w/ FBA w/ FPEA w/ BA w/ PEA w/ PEA:FA (1:1) Fitting linear

21 20

Control w/ FBA w/ FPEA w/ BA w/ PEA

w/ PEA:FA (wt/wt) 1:0 2:1 1:1 1:2 0:1

19 18 17 16

0.5

5

1.0 1.5 Time ( s)

2.0

15

(e)

(f)24

Hysteresis index (%)

-2

12 Control 0

Control Reverse 19.02% Forward 17.32%

-5 -10

w/ PEA:FA (1:1) Reverse 21.15% Forward 20.54%

-15 -20

0.0

(g)

d(J/Jsc) / dV

15

10

5

0 -1.0

0.5 Bias (V)

8

Control w/ FBA w/ FPEA w/ BA w/ PEA w/ PEA:FA (2:1) w/ PEA:FA (1:1) w/ PEA:FA (1:2) w/ PEA:FA (0:1)

FBA

6

PEA: FA (1:1)

FPEA BA

4

12

18.87%

4

0

OC SC

18.31%

Control w/ PEA:FA(1:1)

PEA

2

(h)

20.86 %

20

0

1.0

Slope @ OC

-25

21.30 %

w/ 2D

10

0

50 100 150 200 250 300 Time (sec) OC SC

w/ 2D

0.04 0.03

10

0.02 8

0.01

6 -0.5 V-Voc (V)

0.05

2:1

0.0

Slope @ SC

0.0

PCE (%)

Intensity (a.u.)

10

22

-3

2

(d) Current density (mA cm )

2

FPEA

3

1 4 10

(b)

PEA:FA (1:1) BA PEA

5 Control 4

3

w/ 2D

PCE (%)

-1 -1 2

(c)

7

16

Mobility (cm V S )

(a)

ntrap (×10 cm )

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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1:1

1:2

PEA:FA (wt/wt)

0:1

0.00

Figure 4. Charge recombination losses and Photovoltaic performance and charge collection mechanism. (a) Statistics of trap densities and electron mobilities for the control and 2D/3D films. (b) TRPL spectra of the control and RP/3D films on glass substrates. (c) Statistics of 40-50 devices for the control and 2D/3D perovskite solar cells. (d) J-V curves under both reverse and forward scan directions showing the negligible hysteresis of PEA:FA(1:1)-based 2D/3D devices. (e) Hysteresis index of the control and 2D/3D devices. (f) The stabilized power output of the champion devices measured at a fixed maximum power point (MPP) voltage as a function of


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time. (g) The slope of the normalized J-V curves at each voltage for the control and 2D/3D devices. (h) The slope values at open circuit (OC) and short circuit (SC) for the control and 2D/3D devices. Correlating the changes in device parameters indicates that the enhanced PCE mainly originates from the VOC improvement (Table S1). The high VOC of 1.14 V enables an unprecedented low VOC deficit of 0.40 V, considering a band gap of 1.54 eV for the 2D/3D perovskites. We verified from the film and device characterizations that the built-in band alignment and suppressed non-radiative charge recombination contribute to the improvement of the VOC. Additional consideration is the charge collection at the perovskite/HTL interface. The light J-V curves were normalized (Fig. S12), and the slope as a function of applied voltage was extracted (Fig. 4g).21,49 Fig. 4h illustrates the slope values at the open circuit (OC) and short circuit (SC) for all the films. The charge collection efficiency value at the OC significantly increases from 7.4 for the control 3D device to 9.0, 8.8, 10.1, 10.6 and 12.9 for the FBA-, FPEA-, BA-, PEA- and PEA:FA(1:1)-based 2D/3D devices, respectively. The highest value obtained for the PEA:FA(1:1)-based 2D/3D device explains well its excellent device performance. While the charge collection at the SC remains nearly unchanged (0.01-0.02), indicating similar drift transport for all films because of the presence of a large internal field. These results suggest that the 2D/3D heterojunction with better trap passivation is reflected by the increased charge collection efficiency within the complete devices.


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Ambient stability. Finally, we demonstrated prominently improved ambient stability of the 2D/3D devices. The 2D capping layers with high water resistance act as a protective coating against moisture, preserving the underlying 3D FAPbI3 perovskite (Fig. 5a). The 2D/3D films therefore exhibit stronger water resistance as evidenced by the contact angle analysis (Fig. 5b). The ambient phase stability was investigated for all the films under exposure to air (25 °C, dark, 30-40% relative humidity) for 60 days. Most FA ions are stabilized in the cubo-octahedral cages as T phase for the FBA-, FPEA-, BA- and PEA-based 2D/3D films after ambient exposure (Fig. 5c and 5d). By contrast, conversion from 3C (100) phase and the hexagonal 2H (100) phase is more obvious for the control case under identical conditions. The 2D/3D heterojunction assisted phase stabilization finally leads to improved long-term stability of solar cells. The long-term aging test for non-encapsulated devices was performed under ambient conditions (humidity of 30-40%, dark, 25 °C). The performance was recorded periodically and is shown in Fig. 5e. The 2D/3D devices exhibit significantly enhanced ambient stability compared to the control, particularly for those based on more hydrophobic films. For example the device maintains 84% of its initial value for the FPEA-based film after ambient exposure for 60 days, higher than the PEA:FA(1:1)-based one (52%) despite its lower PCE. While, these values significantly outperform the control one (3%), implying that the 2D/3D hierarchical structure is indeed highly effective not only for stabilizing FAPbI3 in air but also further advancing device performance. Conclusion In summary, we have demonstrated understanding of how interfacial engineering via fluorinating spacer ligands or compositional tuning influences charge transport/extraction and photovoltaic performance of FAPbI3 based 2D/3D solar cells. A templated growth of QWs by the underlying 3D phase was found to be a general property within 2D/3D films irrespective of fluorination or


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compositional engineering. The templated QWs do not follow packing behavior observed within the neat 2D films or single crystals, with more orientation randomness when fluorinating spacers. Interestingly, phase pure n = 2 QWs were achieved for the aliphatic spacer-based 2D/3D films while phase impurity (n = 1, 2) for the aromatic spacers-based films can be improved via compositional engineering. We further confirmed the formation of the built-in band alignment and spontaneous charge transfer from QWs to the underlying 3D layer at the 2D/3D heterojunction. The better QW distribution and corresponding faster charge transfer kinetics enables significantly optimized optoelectronic properties including enhanced carrier mobility, enhanced charge collection, reduced charge non-radiative recombination and reduced loss-inpotential of the 2D/3D devices. The PEA:FA (1:1)-based 2D/3D film achieves a balance between passivation at the 2D/3D heterojunction and charge transport within QWs, leading to a high PCE of 21.15% which is much higher than the 3D counterpart (19.02%) due to significantly improved VOC. Meanwhile, the device exhibits negligible hysteresis and remarkable long-term ambient stability without encapsulation. Such findings provide a deep understanding of the ligandchemistry dependent nature of the 2D/3D heterojunction and delineate precise guidelines for the future design of 2D/3D perovskite solar cells.


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METHODS Materials Preparation 4-Fluorophenethylamin iodide (4-FC6H4CH2CH2NH3I, FPEAI): 84 mL 4-fluorophenethylamin (99%, Acros Organics) and 100 mL hydroiodic acid (45.0%, Sinopharm Chemical Reagent Co., Ltd) solution were mixed and stirred in an ice bath at 0 °C. The reaction solution was stirred for 10 h and dried over vacuum rotary evaporator at 60 °C. The product FPEAI was recrystallized and washed with ethanol and diethyl ether three times until the color turned colorless. Finally, the precipitate was dried at 60 °C in vacuum oven for 12 h. 4, 4, 4 4,4,4

\ \

7 ) 7 )

iodide (CF3(CH2)3NH3I, FBAI) was synthesized by reacting 73 mL with 100 mL hydroiodic acid. The other process is similar with above

demonstration. Solution preparation and device fabrication Solution preparation: The solution preparation was conducted under an inert atmosphere inside a nitrogen glove box. The n-butylamine iodide (BAI, 99.5%, p-OLED) and phenethylammonium Iodide (PEAI, 99.5%, p-OLED) in isopropyl alcohol (IPA, 99.9%, Aladdin) with concentration 5 mg/mL were prepared and stirred overnight at room temperature. The control FAPbI3 solution was prepared with a mixture of 1.2M lead iodide (PbI2, 99.9985%, Alfa Aesar) and 1.2M formamidinium iodide (FAI, 99.5%, p-OLED) dissolved in a mixture of dimethylsulfoxide (DMSO, 99.9%, Sigma-Aldrich), gamma-butyrolactone (GBL, 99%, Sigma-Aldrich) and N,Ndimethylformamide (DMF, 99.8%, Sigma-Aldrich) (volume ratio of 7:7:6) in a glovebox. For the FAPbI3+PbI2 solution, the PbI2 and FAI with the ratio of 1.08:1 were dissolved in the same solvent recipe. The solution was filtered prior to solution-casting. The Spiro-OMeTAD (99.5%,


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p-OLED) solution was prepared by dissolving 90 mg Spiro-OMeTAD and 22 P# lithium bis(trifluoromethanesulfonyl)imide in acetonitrile and 36 P# 4-tert-butylpyridine in 1 mL chlorobenzene. Device fabrication: The FTO-coated glass (2.5 cm × 2.5 cm) was cleaned by sequential sonication in acetone, isopropanol and ethanol for 30 min each and then dried under N2 flow and treated by O3 plasma for 15 min. The TiO2 was prepared by chemical bath deposition with the clean substrate immersed in a TiCl4 (CP, Sinopharm Chemical Reagent Co., Ltd) aqueous solution with the volume ratio of TiCl4:H2O equal to 0.0225:1 at 70 °C for 1 hour. The spincoating was accomplished under inert atmosphere inside a nitrogen glove box. The control FAPbI3 film was deposited using a 1.2 M precursor solution spin-coated at 1000 rpm for 10 s and 4000 rpm for 40 s. At ca. 22 s before the end of the last spin-coating step, 100 P# of chlorobenzene was dropped onto the substrate. The FAPbI3 film was then placed on a hotplate for 30 min at 150 °C. For the 2D/3D film, FAPbI3+PbI2 precursor solution was first spin-coated and annealed following the same procedure as for the control film. The 45 P# iodized ligands in IPA was cast on the thermally-annealed FAPbI3+PbI2 film during spin-coating at 3000 rpm for 30 s, followed by thermal annealing at 100 °C for 3 min. Subsequently, the Spiro-OMeTAD layer was deposited on top of the perovskite by spin-coating at 5000 rpm for 30 s followed by evaporation of 100 nm gold electrodes on top of the stack. Characterization Optical metrology: UV-Visible absorption spectra were acquired on a PerkinElmer UV-Lambda 950 instrument. Steady-state photoluminescence (PL) (excitation at 510 nm) and time-resolved


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photoluminescence (TRPL) (excitation at 510 nm and emission at 806 nm) were measured with a PicoQuant FT-300. First-principles simulations: The plane-wave method in the framework of density-functional theory (DFT) was used as implemented in the VASP code. The projector augmented wave method was used to describe the interaction between ion cores and valence electrons. Spin-orbitcoupling was also considered in the DOS analysis. The van der Waals (vdW) interaction between (PEA)2PbI4 and FAPbI3 counterparts was described by the DFT-D3 correction. A ]

3 k-

point of 3×3×1 for Brillouin zone sampling was generated using the Monkhorst-Pack scheme during calculations. The lattice parameter and atomic positions of all structures were relaxed until the total energy changes became ^ 1.0×10-4 eV. Femtosecond pump-probe transient absorption (TA) measurements were performed at three different power densities (i.e. 30, 70 and 120 P"

2).

The femtosecond laser pulse was

generated by a Ti:sapphire femtosecond regenerative amplifier with 800 nm wavelength and 1 kHz repetition rate (Coherence) and served as both pump and probe beams. The pump pulse with a wavelength of 400 nm and duration of 50 fs generated via a second harmonic generator (SHG) was used to excite all the samples and the probe beam was detected by a high-speed spectrometer (HELIOS, Ultrafast Systems). The wavelength range of the detector was set from 400 to 850 nm, and the spot size of TA is approximately 0.5 mm2 as evaluated by imaging the laser spot. All experiments were carried out at room temperature (i.e. T=300K). Electronic microscopy: The surface morphologies of the perovskite films were characterized by SEM (FE-SEM; SU-8020, Hitachi) at an acceleration voltage of 3 kV and by atomic force microscope (AFM, Dimension ICON).


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X-ray diffraction (XRD) measurements were carried out in a _ _ configuration with a scanning interval of _ between 3 and 45 on a Rigaku Smart Lab (X-ray Source: Cu 1T9

= 1.54 Å).

Grazing incidence wide angle X-ray scattering (GIWAXS) measurements were performed at Dline of the Cornell High Energy Synchrotron Source (CHESS). The wavelength of the X-rays was 1.162 Å with a bandwidth ab b of 1.5%. The scattering signal was collected by a Pilatus 200K detector, with a pixel size of 172 µm by 172 µm placed 173.455 mm away from the sample position. The incidence angle of the X-ray beam was 0.3°. Solar cell characterizations: The J-V performance of the perovskite solar cells was analyzed using a Keithley 2400 SourceMeter under ambient conditions at room temperature, and the illumination intensity was 100 mW cm-2 (AM 1.5G Oriel solar simulator). The scan rate was 0.3 V s-1. The delay time was 10 ms, and the scan step was 0.02 V. The power output of the lamp was calibrated using an NREL-traceable KG5-filtered silicon reference cell. The device area of 0.09 cm2 was defined by a metal aperture to avoid light scattering from the metal electrode into the device during the measurement. The EQE was characterized on a QTest Station 2000ADI system (Crowntech. Inc., USA), and the light source was a 300 W xenon lamp. The monochromatic light intensity for the EQE measurement was calibrated with a reference silicon photodiode. Carrier mobility measurement: Electron-only devices (glass/FTO/c-TiO2/perovskites/PCBM/Ag) were fabricated to measure the electron mobilities of the devices. The dark J-V characteristics of the electron-only devices were measured using a Keithley 2400 SourceMeter. The mobility was extracted by fitting the J-V curves in the space charge limited current (SCLC) regime with the


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Mott-Gurney equation. The trap state density was determined form the trap-filled limit voltage using the equation in the main text. Contact angle measurements were conducted on a DataPhysics OCA 20 instrument with a drop of ultrapure water (0.05 mL). The photographs were taken ca. 1 second after the water was dripped. ASSOCIATED CONTENT Supporting Information is available free of charge on the Internet at http://pubs.acs.org.

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

NOTES The authors declare no competing financial interest.

AUTHOR CONTRIBUTION T. Niu and J. Lu contributed equally to this work. K. Zhao and S.(F.) Liu supervised the work. T. Niu and J. Lu performed most of the experiments. X. Jia, N, Yuan and J. Ding helped with the TA measurements. Z. Xu helped with the first principles simulations. D. Barrit, M. Tang, D. Smilgies and A. Amassian assisted with ex situ GIWAXS measurements. X. Zhang helped with


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the SEM measurements. Q. Li and Z. Liu helped with the trap density measurements. Y. Fan and T. Luo assisted with the TRPL measurements. K. Zhao wrote the manuscript with critical input from all the authors. ACKNOWLEDGMENTS This work was supported by the National Key Research and Development Program of China (2017YFA0204800, 2016YFA0202403), Key projects of the Natural Science Foundation of China (51933010). National Natural Science Foundation of China (61974085), Natural Science Basic Research Plan in Shaanxi Province of China (Program No. 2017JQ6040), the 111 Project (B14041), the National 1000 Talents Plan program (1110010341), and the King Abdullah University for Science and Technology (KAUST). CHESS is supported by the NSF Award DMR-1332208.

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3D heterojunction for high

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