Morphology and Optoelectronic Variations Underlying the Nature of

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Morphology and optoelectronic variations underlying the nature of the electron transport layer in perovskite solar cells Wendy J Nimens, Jonathan Ogle, Anna Caruso, Mackenzie Jonely, Charles Simon, Detlef-M. Smilgies, Rodrigo Noriega, Michael A. Scarpulla, and Luisa Whittaker-Brooks ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.7b00147 • Publication Date (Web): 31 Jan 2018 Downloaded from http://pubs.acs.org on February 5, 2018

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ACS Applied Energy Materials

Morphology

and

Optoelectronic

Variations

Underlying the Nature of the Electron Transport Layer in Perovskite Solar Cells Wendy J. Nimens,† Jonathan Ogle, † Anna Caruso, ‡ Mackenzie Jonely, † Charles Simon, † Detlef Smilgies, # Rodrigo Noriega, † Michael Scarpulla, ‡ and Luisa Whittaker-Brooks†,* †

Department of Chemistry, University of Utah, 315 South 1400 East, Rm 2020, Salt Lake City,

Utah, 84112, USA ‡

Department of Materials Science and Engineering, University of Utah, 315 South 1400 East,

Salt Lake City, UT, 84112, USA #

Cornell High Energy Synchrotron Source, Cornell University, Ithaca, NY 14853, USA

KEYWORDS: perovskite solar cells, electron transport layers, trap states, energy level alignment, titania layers ABSTRACT. Perovskites based on methylammonium lead halides, CH3NH3PbX3 (X = Cl, Br, I) have emerged as one of the most promising materials in solar cell technology. Although the photovoltaics field has witnessed a significant progress in the power conversion efficiency (PCE) of perovskite solar cells, unveiling the contribution of the various factors (i.e., energy level alignment, trap states, electron (hole) mobility, interface interactions, and morphology) affecting

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the observed PCEs is extremely crucial to achieve reproducible and stable devices. This work aims to understand charge transport and recombination within conventional perovskite solar cells due to modifications of the morphology, optoelectronic properties, and energy levels of the titania electron transport layer. Here, we utilize two different processing methods (i.e., solution and sputtering depositions) to yield three morphologically different titania electron transport layers (i.e., planar bulk TiO2, mesoporous TiOx, and sputtered TiO2). We find that the most important factors affecting the PCEs in perovskite solar cells are related to trap-assisted recombination and energy level alignment due to variations in the electron transport layer/perovskite interface. Similarly, we observe that morphology of both the electron transport layer and the perovskite active layer play a minor role on the observed PCEs.

INTRODUCTION With the ever-increasing demand for clean energy, much research effort has been devoted to improving our solar energy technologies. One of the most promising materials in the solar industry

is

the

organic-inorganic

hybrid

perovskites

based

on

the

semiconductor

methylammonium lead halide, CH3NH3PbX3 (X = Cl, Br, I). These haloperovskites when incorporated into solar cells have shown astonishing efficiencies as high as 22.1% and have been considered the biggest energy breakthrough of the past few years.1-5 To date, numerous research investigations have shown that the outstanding power conversion efficiencies attained in perovskite solar cells are attributed to the very intriguing optical and electronic properties of the perovskite layer.6-8 These unique sets of properties include tunable band gaps between 1.2 and 2.3 eV,9, 10 a light absorption spectrum up to a wavelength of 800 nm,7 weak exciton binding energy of about 5-30 meV,11 ambipolar charge transport,12, 13 limited


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charge recombination11, 14, 15 long carrier lifetimes,16 ease of processing, and potential low-cost fabrication.2 Besides exploiting the superb properties observed in organic-inorganic hybrid perovskites, a fundamental understanding of the interactions of the various layers that make up the solar cell and the behavior of charge carriers within the photovoltaic film stack plays an important role for achieving high and reproducible power conversion efficiencies. Although record breaking efficiencies have been obtained with mesoscopic perovskite solar cells, seminal works in the field that have demonstrated the ambipolar charge transport nature of perovskites suggest the possibility of fabricating planar thin film perovskite solar cells with high efficiency and low hysteresis device responses.17,

18

These planar perovskite solar cells are

comprised of either n–i–p or p–i–n device junctions, where n constitutes an electron transport layer (ETL), i is a perovskite layer, and p is a hole transport layer (HTL).19-21 The planar perovskite solar cell architecture discussed in this work is classified as an n–i–p diode junction, where TiO2 and poly[bis(4-phenyl)(2,4,6-trimethylphenyl)amine], PTAA ̶ a doped small molecule ̶ are used as the ETL and HTL, respectively. Specifically, the ETL as a vital component of a perovskite solar cell should be engineered and optimized to reduce losses stemming from poor charge transport and selectivity, charge recombination, conduction band misalignment, and build-in potential. Such engineering and optimization of the ETL include the synthesis of new ETLs, fabrication of dense and pinhole─free ETLs, and an efficient charge extraction via a suitable energy level alignment with the perovskite absorber. TiO2 has been demonstrated as one of the most efficient ETL in perovskite solar cells due to its high transparency, physical stability, fast electron injection rates, long electron lifetimes, and


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proper energy level alignment with perovskites (particularly, methylammonium lead iodide, CH3NH3PbI3).22, 23 However, there are major drawbacks often observed when using TiO2 as the ETL in planar perovskite solar cells. For example, energetically deep trap states in TiO2 are readily formed upon UV light exposure.24, 25 These trap sites act as recombination centers that detrimentally impact the short-circuit current densities and fill factors of solar cell devices. Moreover, defects at the TiO2/perovskite interface are key contributors to the hysteresis effects often observed in perovskite solar cells thus strongly suggesting the need for ETLs with fewer trap states.26 Additional challenges to overcome in TiO2 are its low conductivity and electron mobility compared to those of other materials. The incorporation of dopants (viz. Li+,27 Nb4+,28 Zn2+,29 Y3+,29 Zr4+,29 and Mo4+,29) within the TiO2 crystal structure has been demonstrated as an effective route for increasing its conductivity from 1.42 x 10-4 to 1.86-3.72 x 10-4 S cm-1.29 While effective for increasing the conductivity of TiO2, dopant incorporation within the TiO2 crystal structure requires an additional processing step during synthesis and often utilizes rare and expensive reagents. In a similar vein, metal oxides with higher electron mobilities (10-300 cm2 V-1s-1 vs 0.1-4 cm2 V-1s-1 for TiO2), such as ZnO,22, 30 SnO2,22 and WOx,22 have successfully been incorporated into planar perovskite solar cells.

However, issues with degradation,

hysteresis, and energy level alignment in these materials have caused TiO2 to remain the most commonly used ETL in perovskite solar cells.22,

31, 32

Thus, investigating and optimizing the

properties of undoped TiO2 can allow a maximum power conversion efficiency to be reached while retaining the ease of fabrication and low cost that is typically associated with the use of TiO2 in solar devices. Another drawback is that the solar cell performance of devices having TiO2 as the ETL has been proven to be strongly dependent on the details pertaining the fabrication methods of TiO2 thin


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films. Various methods such as spin-coating,29,

33, 34

spray pyrolysis,35-37 sputtering,38-40 dip-

coating,41 and atomic layer deposition42-45 have been used to fabricate TiO2 thin films. These fabrication methods have yielded TiO2 thin films with irreproducible and hard-to-control optoelectronic properties. In this work, we investigated the morphology, defects, and electronic contributions of TiO2 as an ETL in chlorine-doped methylammonium lead iodide (CH3NH3PbI3xClx)

perovskite solar cells. We compared three different processing conditions for TiO2; bulk

solution-processed (b-TiO2), mesoporous solution-processed (m-TiOx), and sputtered (sp-TiO2). Furthermore, we studied the effects of structure, crystallinity, charge transport, energy level alignment, and trap-state density in the TiO2 ETL on the efficiency of perovskite solar cells. Our studies not only provide a detailed understanding on how fabrication and optimization of TiO2 thin films can be controlled towards an efficient electron collection and hole blocking layer but also offer significant insights into the underlying physical and electronic nature affecting the performance of TiO2 when incorporated into perovskite solar cells. RESULTS AND DISCUSSION The ETL morphology, structure, and optoelectronic properties have been proven to strongly impact the short circuit current (Jsc), open-voltage circuit (Voc), and fill factors (FF) in perovskite solar cells.22,

23

Thus, the ETL in perovskite solar cells should be fabricated to achieve low

surface roughness, pinhole-free and highly uniform films, low resistance, low trap densities, and optimal energy level alignment. In order to control and investigate these various properties, three different ETLs comprised of TiO2 were fabricated. Two of these TiO2 ETLs were solution processed to yield compact TiO2 (b-TiO2) and mesoporous TiOx (m-TiOx) films whereas, a third TiO2 ETL was sputter deposited atop a substrate. The TiO2 ETLs’ structure and morphology were thoroughly studied. X-ray diffraction measurements conducted on the various ETLs reveal


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structural and crystalline differences associated with the TiO2 crystal structure. Figure S1 displays the XRD patterns for the ETLs. Consistent with previously reported XRD studies on mesoporous titania layers,46 no diffraction peaks are observed in the XRD spectrum of our asprepared m-TiOx thin films suggesting the formation of an amorphous TiOx layer. It is worth noting that the m-TiOx thin film was calcined at 450 oC to remove the surfactant used in its fabrication.

However, this high annealing temperature did not induce the formation of a

crystalline layer hence remaining vastly amorphous.

In contrast, the XRD spectra for b-TiO2

and sp-TiO2 show diffraction peaks at 25.4o and 38.1o further confirming the crystalline nature of these ETLs. These diffraction peaks are associated with the (101) and (004) crystallographic planes of the anatase structure of TiO2, respectively. Atomic force microscopy (AFM) was used to investigate the surface morphology of the ETLs as a function of processing conditions. Figure 1A shows the AFM image of a 30 nm-thick b-TiO2 thin film deposited on an ITO-coated glass substrate. The average root mean squared (RMS) surface roughness is 2.3 nm. The AFM image of a 30nm–thick m-TiOx thin film is presented in Figure 1B. The average RMS surface roughness for this film is 1.4 nm, suggesting the formation of a smoother film compared to the bTiO2 thin film. Also, the calcination of the surfactant used in the synthesis of m-TiOx leaves behind disordered, spherical pores within an interconnected network of TiOx, which appear as circular dark regions in the AFM images. This process imparts the mesoporous nature that is characteristic of our m-TiOx layers. Compared to the solution processed films, the sp-TiO2 thin film is the smoothest; the RMS surface roughness of the sp-TiO2 layer is 0.5 nm (Figure 1C). As reference, the RMS surface roughness of an ITO-coated glass substrate is 3.8 nm; the deposition of these ETLs thus appears to planarize the ITO surface. This feature is extremely important since it suppresses any pinholes and sharp spikes that may be present in ITO and may


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contribute to a higher power conversion efficiency when these smooth layers are incorporated into perovskite solar cells (vide infra).47

Figure 1. AFM topography images for A) b-TiO2, B) m-TiOx, and C) sp-TiO2 electron transport layers.


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We turn now to the determination of the optical properties and defect concentration of our ETLs as a function of processing conditions. Figure 2A displays the UV-vis-NIR absorption spectra for b-TiO2, m-TiOx, and sp-TiO2 ETLs. The absorbance versus wavelength curves are used to estimate the optical bandgap of our ETLs as a function of processing conditions. The absorbance onset in b-TiO2, m-TiOx, and sp-TiO2 ETLs appears at 360, 355, and 381 nm, respectively corresponding to optical bandgaps of 3.4, 3.5, and 3.3 eV, respectively. We believe the variations in the optical bandgaps may stem from morphology and uniformity variations of the various ETLs upon different processing conditions. To investigate the presence of defects in our ETLs, photoluminescence (PL) spectroscopy studies were carried out. Figure 2B shows the PL spectra for 30nm-thick b-TiO2, m-TiOx, and sp-TiO2 thin films. Two distinctive features are evident in the PL spectra of our TiO2 thin films. The first feature is centered at ≈ 368 nm in the ultraviolet region and it is dubbed the near band-edge emission due to the electronic transition of electrons from the conduction band (CB) into the valence band (VB).48, 49 The second feature spans from 390-550 nm in the visible region, and it corresponds to both bulk and surface defects in TiO2.48, 49 Such defects may be related to the existence of singly ionized oxygen vacancies, recombination of a photogenerated hole with an electron occupying an oxygen vacancy, impurities or defects in the TiO2 crystal structure, and/or titanium atoms lying in the interstitials.48, 49 As shown in Figure 2B, we observe the evolution of a broad peak ranging between 390-560 nm for b-TiO2 and m-TiOx. We believe the solution processes employed in the fabrication of b-TiO2 and m-TiOx due to the use of surfactants trigger the formation of defects within the thin films. The heightened defect concentration of m-TiOx is likely due to its increased surface area, allowing more opportunity for surface defect formation


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when compared to the non-porous b-TiO2 surfaces. On the other hand, surface defects are substantially diminished in the sp-TiO2 thin film. The fact that a TiO2 sputter target is the only precursor used in the fabrication of sp-TiO2 thin films helps to inhibit the introduction of undesired intrinsic and extrinsic defects within the TiO2 crystal structure.

Figure 2. A) UV-vis-NIR and B) PL spectra for as-fabricated TiO2 thin films.


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To further investigate defect concentrations quantitatively, we fabricated capacitors based on our ETLs and measured their capacitance-voltage (C-V) characteristics in the dark. When a bias is applied across an ETL device (ITO/ETL/Ag), charges reorganize and a capacitance is built over the ETL. When the bias is reversed, the stored charge changes and produces an S-shaped C-V curve as shown in Figure 3.50 If defect states are present in the ETL, charges will get trapped and hysteresis in the C-V will develop. The magnitude of the hysteresis loop will strongly depend on the defect density present in the ETL. Hence, by analyzing the magnitude of C-V hysteresis loop, we calculate the trap density (∆Not) for each our ETLs as per equation 1:

Here, Cox is the capacitance of the ETL under accumulation, ∆Vmg is the voltage difference obtained from the C-V hysteresis, q is the charge of an electron and A is the device area. Quantifying the ∆Vmg in our C-V data allows us to estimate defect concentrations in our ETLs. From Figure 3A, we determine a ∆Vmg for the b-TiO2 ETL to be 70 mV, which results in a defect concentration of 2.6 x 1014 m-2. The ∆Vmg for the capacitor fabricated with a m-TiOx ETL is 210 mV, resulting in a defect density of 3.1 x 1014 m-2 (Figure 3B). Although both b-TiO2 and m-TiOx ETLs are solution processed, we observe that the defect concentration in m-TiOx is 1.2 times higher than that of b-TiO2. As mentioned above, we attribute the higher trap density in mTiOx to its increased surface area and surface trap states compared to that of b-TiO2. Since surface area contact between the ETL and the perovskite layer plays a significant role in charge


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transport, we would expect solar cell devices with m-TiOx to perform better than b-TiO2 in spite of the higher trap density. As an interesting note, we observe that the S-shaped C-V curve in bTiO2 is inverted in contrast to that of m-TiOx. We believe this discrepancy may originate from a larger interplay between surface and bulk defects in b-TiO2 than that of m-TiOx. Moreover, as evidenced in Figure 3C, the capacitor based on the sp-TiO2 ETL shows a ∆Vmg of 3 mV, which corresponds to a defect density of 8.6 x 1012 m-2. The very low trap density in sp-TiO2 indicates that devices with this ETL should experience fewer losses due to trap-assisted recombination and exhibit a higher PCE ─when incorporated into solar cells─ than the solution processed ETLs (vide supra).


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Figure 3. C-V characteristics of capacitors based on A) b-TiO2, B) m-TiOx, and C) sp-TiO2 electron transport layers (ETLs).

The magnitude of the hysteresis loop observed in the C-V

spectra is proportional to the defect concentration in our ETLs.


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Next, we investigated how the different physical and chemical properties observed in the ETLs impact the overlying perovskite layer.

Figure 4 displays SEM micrographs of the bulk

crystalline perovskite (CH3NH3PbI3-xClx) atop each of the ETLs. An SEM image of a perovskite layer atop a bare ITO-coated glass substrate is presented as the control (Figure 4A). We observe that a compact film with homogeneous grain size distribution (average value of 1.6 ± 0.1 µm) is obtained for a perovskite layer grown atop bare ITO.

Moreover, we do not observe any

significant variation in the morphology of the CH3NH3PbI3-xClx layer when it is deposited atop b-TiO2 or m-TiOx ITO-coated glass substrates (Figure 4B and 4C). Due to the relatively thin m-TiOx layer (30-nm thick), the perovskite layer (500-nm thick) is likely to infiltrate the ETL as well as form a capping layer atop the substrate,51 thus it is not surprising to see that there are subtle variations in the morphology of the perovskite layer when cast onto m-TiOx or b-TiO2 ITO-coated glass substrates. However, for the CH3NH3PbI3-xClx layer grown on the sp-TiO2 ITO-coated glass substrate (Figure 4D), we observe the formation of a compact film with slightly smaller grain sizes (0.7 ± 0.1 µm).

In addition, we do not observe the formation of

pinholes as we alter the processing conditions of the underlying ETL layer.

Further

investigation of the morphology, structure, and orientation of the CH3NH3PbI3-xClx layer upon interaction with the ETLs were performed via grazing incidence wide angle X-ray scattering (GIWAXS) studies. Figure 5 shows indexed 2D GIWAXS images of CH3NH3PbI3-xClx thin films atop bare ITO and the ETL deposited on ITO-coated glass substrates (variations in the breadth of the peaks in the different GIWAXS profiles arise from differences in the sizes of the substrates measured).

Here, the GIWAXS profiles display the formation of a crystalline

CH3NH3PbI3-xClx layer. GIWAXS scattering peak assignments corroborate the formation of a


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tetragonal CH3NH3PbI3-xClx structure with its most intense scattering peaks corresponding to reflections associated with the (110) and (220) crystallographic planes at q = 1.04 A-1 and q = 2.08 A-1, respectively.52

The GIWAXS patterns presented in Figure 5 also demonstrate

anisotropic intensities around the (110) and (220) family of planes. In particular, the (110) reflection of CH3NH3PbI3-xClx is most intense at the meridian (qxy = 0), an indication that the (hk0) planes of CH3NH3PbI3-xClx are preferentially oriented parallel to the substrate.

Figure 4. Top view SEM micrographs of CH3NH3PbI3-xClx thin films atop A) an ITO-coated glass substrate, B) b-TiO2 on an ITO-coated glass substrate, C) m-TiO2 on an ITO-coated glass substrate, and D) sp-TiO2 on an ITO-coated glass substrate. Scale bar: 10 µm


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Figure 5. GIWAXS images of CH3NH3PbI3-xClx grown atop an A) ITO-coated glass substrate, B) b-TiO2 on an ITO-coated glass substrate, C) m-TiO2 on an ITO-glass coated substrate, and D) sp-TiO2 on an ITO-coated glass substrate. For clarity, we have labelled the reflections associated with CH3NH3PbI3-xClx crystallographic planes in (A). Samples underwent degradation during shipping as evidenced by the evolution of a scattering peak associated with the precursor PbI2. Scattering peaks at q values < 0.8 A-1 are also observed upon degradation of the sample. These peaks correspond to the formation of an intermediate phase upon water intercalation within the PbI2 lattice. Peak broadening is due to differences in sample sizes and not to changes in crystallite sizes.


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An alternative means of quantifying the extent by which the ETLs affect the orientation of the deposited CH3NH3PbI3-xClx layer is to estimate the Hermans’ orientation function about the (110) reflection of the CH3NH3PbI3-xClx layer, f110:53

Where,

and Φ is the azimuthal angle; I(Φ) is the intensity of the (110) reflection as a function of the azimuthal angle. Accordingly, f110 can range from 1 to -0.5 wherein f110 = 1 indicates perfect alignment of the CH3NH3PbI3-xClx layer normal to the substrate; f110 = -0.5 indicates perfect alignment of the CH3NH3PbI3-xClx layer in the plane of the substrate; and f110 = 0 indicates no preferential alignment of the CH3NH3PbI3-xClx layer. Per equations 2 and 3, we determine f110 = 0.63, 0.68, 0.72, and 0.62 for the CH3NH3PbI3-xClx layer deposited atop bare ITO, b-TiO2, mTiOx, and sp-TiOx, respectively. The calculated f110 displays only small differences amongst all substrates thus we conclude that the deposition method of these ETLs does not significantly influence the orientation nor the morphology of the CH3NH3PbI3-xClx layer. Despite the absence of dramatic crystallographic and morphological differences amongst CH3NH3PbI3-xClx films deposited on the different ETL layers, the interplay between the physical and electronic structure of different ETL/CH3NH3PbI3-xClx interfaces may have a noticeable influence on the optoelectronic properties of CH3NH3PbI3-xClx thin films, especially when


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incorporated into solar cells.

Thus, elucidating the electronic level alignment of the

ETL/CH3NH3PbI3-xClx hybrid is of paramount importance for efficient charge extraction (injection) and open-voltage circuit (Voc) loss minimization. Using ultraviolet photoemission spectroscopy (UPS), we probed the energetics of all pristine layers (viz., b-TiO2, m-TiO2, spTiOx, CH3NH3PbI3-xClx) and that of a 500–nm thick CH3NH3PbI3-xClx layer grown on top of the three TiO2 ETL types. UPS measurements on a CH3NH3PbI3-xClx layer grown on top of an ITOcoated glass substrate were also performed as a control experiment. We are aware that ion migration within the CH3NH3PbI3-xClx layer as well as band bending in both the CH3NH3PbI3xClx

layer and our ETLs can strongly complicate the energetics in our devices.54,55 Photoemission

spectroscopy studies investigating these two major factors in perovskite solar cells are currently underway. Also, all samples were handled with extreme precaution in an inert environment to minimize the impact of surface contaminants on the measured electronic structure. Figure S2 displays the direct photoemission spectra of pristine b-TiO2, m-TiO2, and sp-TiOx ETLs where all energy levels are referenced to a common Fermi level (0 V). Here, the high binding energy region of the UPS spectra is used to determine the secondary electron cutoff energy (ESECO). The ESECO is subtracted from the He I excitation energy (21.2 eV) to determine the work function (WF) of the ETLs (Figure S2A). The WF of pristine b-TiO2, m-TiO2, and sp-TiOx ETLs is 3.6, 3.7, and 3.5 eV, respectively. Figure S2B presents the UPS spectra of the highest lying valence band regions. Here, the valence band edge (VB) is determined by linear extrapolation of the valence band onset subtracted to the background around the Fermi level. The VB of pristine bTiO2, m-TiO2, and sp-TiOx ETLs is 3.2, 3.1, and 3.2 eV below the Fermi level, respectively. Furthermore, we determined the ionization energies (IE) to be 6.8 ± 0.1, 6.8 ± 0.1, and 6.7 ± 0.1 eV for our b-TiO2, m-TiO2, and sp-TiOx, ETLs, respectively. These calculated IEs are in close


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agreement with reported values.56,

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While the position of the VB has been determined

predominantly by UPS measurements, the position of the conduction band edge (CB) relative to the vacuum level (electron affinity, EA) has been ordinarily obtained by subtracting the optical bandgap from the VB. By performing this exercise, a CB minimum (CBM) at 0.2, 0.4, and 0.1 eV above the Fermi level is obtained for our b-TiO2, m-TiO2, and sp-TiOx ETLs, respectively. This leads to electron affinities of 3.4 ± 0.1, 3.3 ± 0.1, and 3.4 ± 0.1 eV, respectively. Although we observe subtle differences in the VB onsets, IE, and EA of the various titania layers upon different defects concentrations, we do observe substantial variations in the WF (particularly, when going from a m-TiOx to a sp-TiO2 ETL) that ultimately affect energy level alignment when these layers are incorporated into solar cells. These WF differences may be related to variations in film morphologies and defect concentrations of the various ETLs undergoing different processing conditions. Similar variations in the electronic properties upon different processing conditions of phenyl-C61-butyric acid methyl ester (PCBM) when used as an ETL in perovskite solar cells were recently reported, thus pointing out the importance of elucidating the influence of structural order and energy disorder as a function of processing conditions in other ETL systems.58 We then investigated the effects that the deposition of the different ETLs have on the electronic structure of the CH3NH3PbI3-xClx layer. UPS studies have shown that the interplay between different layers in the solar cell can have significant effects on the electronic properties of a material. Specifically, the use of an ETL with different electronic properties provide different doping and charge transfer characteristics at the ETL/perovskite interface, leading to different VBM and work function values.57, 59,

60

Figure S3 shows the direct photoemission spectra of

CH3NH3PbI3-xClx atop the b-TiO2, m-TiO2, and sp-TiOx ETLs. The WFs of the CH3NH3PbI3-


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xClx

layers are 4.2 ± 0.1, 4.0 ± 0.1, and 4.1 ± 0.1 eV when spin-coated on top of b-TiO2, m-TiO2,

and sp-TiOx ETLs, respectively (Figure S3A).

The position of the CH3NH3PbI3-xClx VB

maximum is at 1.4, 1.5, and 1.5 eV below the Fermi level for samples deposited on b-TiO2, mTiO2, and sp-TiOx ETLs, respectively (Figure S3B). Interestingly, the IE centered about 5.6 ± 0.1 eV for the CH3NH3PbI3-xClx remains unchanged, indicating that the deposition of the ETLs does not induce a significant interface dipole. Similarly, CB minima at 0.2, 0.1, and 0.1 eV above the Fermi level are obtained, leading to EAs of 4.0 ± 0.1, 3.9 ± 0.1, and 4.0 ± 0.1 eV for samples deposited on b-TiO2, m-TiO2, and sp-TiOx ETLs, respectively. Also, it is worth noting that the IE and EA values for the perovskite layer calculated in this work are slightly but significantly different from previous reports (IE = 5.4 eV and EA = 3.7 eV) underlying the importance of the ETL on the electronic properties of the CH3NH3PbI3-xClx layer.56, 57 Finally, by combining the UPS results along with the bandgap values obtained from our UV-visNIR absorption measurements, we were able to construct the heterojunction energy level diagrams displayed in Figure 6. We observe the formation of type-II heterojunctions for devices having b-TiO2 and sp-TiO2 as the electron transport layers. This result is in agreement with previous experimental results and DFT calculations,61 which promote electron transfer from the CH3NH3PbI3-xClx layer to the ETLs. Furthermore, we observe differences in the alignment of the CBM position of the CH3NH3PbI3-xClx layer with the CBM of the various ETLs (∆CBM = CBMCH3NH3PbI3-xClx ─ CBMETL). As shown in Figure 6, ∆CBM offsets of 0.02, 0.3, and 0.01 eV are obtained when b-TiO2, m-TiO2, and sp-TiOx thin films are used as the ETLs, respectively. These differences in energy level offsets may be due to varying compositions of the TiO2 layer, contaminants at the interface, trapped charges at the interfaces, and/or trapped charges in the bulk of TiO2.

Moreover, these results imply that a heterojunction based on sp-


19


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TiO2/CH3NH3PbI3-xClx hybrid layers should yield a better interface for electron extraction from the CH3NH3PbI3-xClx absorber layer into the sp-TiO2 electron transport layer. In addition, the relative energy differences between the CBM of the ETLs and the VBM of CH3NH3PbI3-xClx can give us an estimate of the expected trend in Voc. In Figure 6, this difference decreases from spTiO2 to b-TiO2 to m-TiOx. From these energy level alignments, we expect to achieve the highest Voc from a heterojunction having sp-TiO2 as the ETL, followed by b-TiO2, then m-TiOx.

Figure 6. Energy level diagrams for pristine ETLs and CH3NH3PbI3-xClx films deposited on top of our ETLs.

These energy level diagrams were derived from UPS and UV-vis-NIR

measurements. For clarity, we have provided the CH3NH3PbI3-xClx bandgap (Eg), Fermi level


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position (EF), electron affinity (EA), ionization energy (IE), and work function (WF) of the various layers.

Shifting to the fabrication and testing of perovskite solar cells, in Figure 7 we present the device characteristics of planar heterojunctions having CH3NH3PbI3-xClx as the absorber layer and bTiO2, m-TiO2, and sp-TiOx as the ETLs. We fabricated and tested solar cells with a conventional architecture

of

ITO/ETLs

(30nm)/CH3NH3PbI3-xClx(600nm)/PTAA(200

nm)/Au(100nm).

Operational testing with an aperture defining the illuminated area minimized potential overestimation of the photocurrents in our devices and thus reflecting the accurate determination of the power conversion efficiency of individual devices.

The current-voltage (J-V)

characteristics of devices having b-TiO2, m-TiO2, and sp-TiOx as the ETLs under illumination are shown in Figure 7A (the complete set of device characteristics using these ETLs is tabulated in Table S1). Of 12 devices tested across three chips using b-TiO2 as the ETL, average Voc, Jsc, and FF of 0.77 ± 0.03 V, 18.9 ± 1.5 mA cm-2, and 41 ± 2 % were obtained, respectively. These values lead to an average power conversion efficiency (PCE) of 6.0 ± 0.3 %. Likewise, when solar cells are fabricated with m-TiOx as the ETL, a Voc, Jsc, and FF of 0.76 ± 0.02 V, 19.7 ± 1.3 mA cm-2, and 49 ± 3 % are obtained, respectively, corresponding to a PCE of 7.3 ± 0.4 %. Solar cells using sp-TiO2 as the ETL have enhanced device characteristics; Voc, Jsc, and FF of 0.88 ± 0.01 V, 23.1 ± 0.8 mA cm-2, and 64 ± 1 %, respectively, yield an average PCE of 13.2 ± 0.3 %. We attribute the higher performance of sp-TiO2 primarily to the reduced trap density in this ETL and its superior energy level alignment with respect to the CH3NH3PbI3-xClx layer. In addition, the smooth surface fabricated during the sputter deposition process provides a pristine interface between the sp-TiO2 and the CH3NH3PbI3-xClx layer, contributing to an increase in the FF. With


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a minimal trap density and very few defects, trap-assisted recombination –as discussed later- is kept to a minimum and charge transport is maximized thus contributing to a higher Jsc value when compared to devices having b-TiO2 and m-TiOx as the ETLs. While the trap-density in mTiOx thin films is higher than that of b-TiO2 thin films and a less favorable energy level alignment with the CH3NH3PbI3-xClx absorber is observed, the superior performance of solar cells having m-TiOx as the ETL when compared to solar cells with b-TiO2 ETLs stems from its high surface area architecture, which facilitates charge transport and aids in its increased Jsc relative to b-TiO2. The increased surface area enhances the intimate contact of the m-TiOx with the CH3NH3PbI3-xClx and ultimately decreases the contact resistance of the device, overcoming the detrimental effects of trap-assisted recombination in m-TiOx. We point out that these devices suffer from hysteresis in J-V (Figure S4). This hysteresis (which is deeper for perovskite cells fabricated in a conventional architecture) has been demonstrated to dramatically alter J-V characteristics and impact PCE values.26, 62-65 To quantify the hysteresis effects in our solar cells as a function of different ETLs, we determined the dimensionless hysteresis index (HI) as per equation 4:66, 67

Where, J(reverse)(Voc/2) and J(forward)

(Voc/2)

are the photocurrent at 50% of Voc for the reverse and

forward scans, respectively. Forward and reverse scan mean scanning from Jsc to Voc and vice versa, respectively. Accordingly, HI can range from 0 to 1, wherein a HI of 0 corresponds to a cell with negligible hysteresis, while a HI of 1 corresponds to a device with a hysteresis contribution as high as the Jsc.

A HI of 0.43, 0.36, and 0.20 are obtained for solar cells


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comprised of b-TiO2, m-TiO2, and sp-TiOx as the electron transport layers.

Traditionally, it

would be expected that a m-TiOx scaffold would yield a device with the lowest HI, however due to the presence of a thick perovskite capping layer atop the m-TiOx, we do not expect the HI to be minimized.2 This suggests that hysteresis effects in CH3NH3PbI3-xClx conventional solar cells are minimized when sp-TiO2 is used as the ETL. Although, we have observed very large discrepancies in PCE performances between the forward (-0.2 V to 1 V) and reverse (1 V to -0.2 V) scans, the underlying mechanism on why this happens in our cells has remained elusive so far. We believe such discrepancies may be related to a slow dynamic process of capacitive current,63, 66-68 slow ion migration,26, 66, 67, 69-71 trapping/de-trapping of charge carriers,26, 72 and/or ferroelectricity.26, 62, 64, 71, 73-76 Further verification that our photocurrents are those of individual cells stem from external quantum efficiency (EQE) measurements. We have included in Figure 7B the EQE spectra for the same devices whose J-V characteristics are shown in Figure 7A. The devices show a broad EQE spectrum spanning a wavelength range of 300-850 nm; photocurrent generation at these wavelengths is attributed to photoabsorption and current collection by the CH3NH3PbI3-xClx layer. Here, the EQE results at short-circuit conditions yield an integrated Jsc of 20.5, 20.8, and 22.7 mA cm-2 for solar cells with b-TiO2, m-TiOx and sp-TiO2 ETLs, respectively. The superior EQE results of the sp-TiO2 device arise from the higher current output of the device, relative to the solution processed ETLs. These values are in close agreement with the measured Jsc extracted from the J-V characteristics, which is an indication of the absence of any parasitic leakage currents. To investigate the effect that the various ETLs has on the dominant recombination mechanism in our devices, J-V characteristics were monitored as a function of light intensity. Figure 7C


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displays the power law dependence of the Jsc under different light intensities ranging from 10 to 110 mW cm-2. A power law relationship with an exponent α = 0.75 is expected for a solar device that is space charge limited due to a carrier imbalance or an interfacial barrier, while a power law relationship with an α value close to 1 suggest no space charge effects.77-79 All of our devices show α values close to 1, suggesting negligible energy barrier due to bimolecular recombination effects. While bimolecular recombination appears to be low in these devices, we believe that trap-assisted recombination provides the main pathway for carrier recombination. A simplified Shockley Reed Hall mechanism can be used to determine the trap assisted recombination effects in solar cells by examining the relationship between Voc and the log of the light intensity. Since there is no current extraction occurring at the Voc, upon excitation, all carriers must recombine either bimolecularly or via trap-assisted recombination. The changes in Voc with light intensity are therefore very sensitive to the presence, or lack, of trap states in the device. The change in Voc (δVoc) as a function of light intensity has been shown to linearly scale with (kTq-1). Where, k is the Boltzmann constant, ܶ is temperature, and q is the charge of an electron. If the slope exceeds (kTq-1), energy loss due to trap-assisted recombination is present.80, 81 The slopes of the Voc versus the log of the light intensity are shown in Figure 7D. From the relationship between Voc ~ log (light), a slope of ≈1.10kTq-1 is determined for both devices comprised of b-TiO2 and m-TiOx and the ETLs, indicating that trap-assisted recombination is present near Voc. A slope of ≈ 0.44kTq-1 is obtained for solar cells with sp-TiO2 ETLs, suggesting that trap-assisted recombination in these devices is minimal. To further examine charge recombination in our devices, the dark J-V characteristics were fit to the ideal diode model, following the Shockley equation:80, 82, 83


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Where, Jph is the photogenerated current, J0 is the dark saturation current, n is the diode ideality factor, Rs is the series resistance, and Rsh is the shunt resistance. Dark curves for our solar cells and the corresponding fits are presented in Figure S5. As J0 is increased dramatically by non– radiative recombination, a lower J0 value of 1.95 x 10-13 mA cm-2 for CH3NH3PbI3-xClx solar cells with sp-TiO2 ETLs compared to those of CH3NH3PbI3-xClx solar cells having b-TiOx (J0 = 1.44 x 10-11 mA cm-2) and m-TiO2 (J0 = 6.87 x 10-8 mA cm-2) ETLs reveals lower carrier recombination and trap states either at the TiO2/CH3NH3PbI3-xClx interfaces or within the CH3NH3PbI3-xClx bulk layer. The calculated Rs (21 Ω) is lower in the solar cells having sp-TiO2 as the ETL than those in cells comprised of b-TiO2 (Rs = 40 Ω) and m-TiOx (55 Ω) ETLs. These results indicate that carrier transport is more favorable in CH3NH3PbI3-xClx solar cells comprising sp-TiO2 ETLs. In a similar vein, if trap-assisted recombination is present, n will be close to 2. n values for solar cells having b-TiO2, m-TiOx and sp-TiO2 ETLs are found to be 1.9, 1.8 and 1.5, respectively. These results are consistent with the light dependence study, finding that trap-assisted recombination is most prevalent in b-TiO2 and m-TiOx than in sp-TiO2. The lower n value for the sp-TiO2 device is representative of its low trap density. In toto, the reduced carrier recombination and lowered series resistances in solar cells comprised of sp-TiO2 ETLs are in agreement with the higher Jsc and FF characteristics observed in the illuminated devices.


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Figure 7. Device characteristics of CH3NH3PbI3-xClx solar cells. A) J-V curves for solar cells of CH3NH3PbI3-xClx deposited atop the various ETLs. Inset shows the device architecture. B) EQE spectrum as a function of wavelength of monochromatic irradiation for the device whose J–V characteristics are shown in (A). C) Jsc versus light intensity. D) Voc versus light intensity.

Electrical impedance spectroscopy (EIS) is a complementary technique to ideal diode analysis of the J-V data yielding insights into the charge transport processes and resistances. Figure 8 presents EIS data for the three types of devices. When fit to the equivalent circuit also presented in Figure 8, the EIS data allows for the determination of the contact (Rco) and recombination (Rrec) resistances of our devices. The Rco of each device relates to the resistance at the ETL/


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CH3NH3PbI3-xClx interface.81 A lower Rco indicates more efficient electron transport across the ETL/CH3NH3PbI3-xClx interface. Rrec is the resistance to recombination, thus larger Rrec improves PCE in a device by preventing recombination losses.81 As shown in Figure 8, the Nyquist plots for CH3NH3PbI3-xClx solar cells with b-TiO2, m-TiOx, and sp-TiO2 biased at the Voc under dark conditions demonstrate distinctive resistance behaviors where the main arc is attributed to charge extraction from the CH3NH3PbI3-xClx into the ETLs.84, 85 Here, a reduction in Rco from 306 to 73 Ω is obtained for cells having b-TiO2 and m-TiOx as ETLs, respectively. This reduction in Rco is likely due to the increased surface area contact and charge transport across the interface between the CH3NH3PbI3-xClx and m-TiOx. An intermediate Rco of 177 Ω is found in devices having spTiO2 as the ETL. We believe this is due to the ETL having an optimal interface with the perovskite layer but reduced surface area relative to that of m-TiOx ETLs. The lower Rrec values found in CH3NH3PbI3-xClx solar cells with m-TiOx (3337 Ω) and b-TiO2 (3994 Ω) ETLs, suggest increased losses due to trap-assisted recombination when compared to that of cells with sp-TiO2 (4976 Ω) ETLs. This is in agreement with our defect studies acquired from PL and C-V measurements. The results from the EIS analysis indicate that Rrec and Rco have large impacts on device performance. Here, sp-TiO2 has the most resistance to trap-assisted recombination while maintaining a reduced Rco that is optimal for a high PCE. Although b-TiO2 has a slightly higher Rrec than that of m-TiOx, its elevated Rco results in a less than ideal electron transport layer for CH3NH3PbI3-xClx solar cells which is further corroborated by its lower PCEs.


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Figure 8. Electrochemical impedance spectroscopy (EIS) characteristics for solar cell devices with b-TiO2, m-TiOx, and sp-TiO2 as the ETLs. Here, Nyquist plots were acquired in a range of 500 Hz to 5 MHz at bias corresponding to the Voc of the J-V curve.

To investigate the effect of traps (interfacial and bulk) on charge-carrier dynamics in CH3NH3PbI3-xClx

solar

cells

with

the

various

ETLs,

we

employed

time-resolved

photoluminescence (TRPL). TRPL traces were measured for three nominally identical samples for each of the following preparations, i.e., CH3NH3PbI3-xClx on bare glass, CH3NH3PbI3-xClx on bare ITO, and CH3NH3PbI3-xClx atop each of the ETLs (b-TiO2, m-TiOx, and sp-TiO2). Each decay was fit to a single- or bi-exponential decay convoluted with the instrument response function (IRF). The results for the average parameters for each preparation are shown in Table 1.


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Table 1. Average decays and relative amplitude for CH3NH3PbI3-xClx deposited on b-TiO2, mTiOx, and sp-TiO2, respectively. TRPL decays and amplitudes are also presented for a CH3NH3PbI3-xClx layer deposited on bare glass and bare ITO. Sample

Average τ1

Glass / CH3NH3PbI3-xClx

1.03 ± 0.01 ns

Average Relative Amplitude 0.27 ± 0.01

Average τ2

3.05 ± 0.06 ns

ITO / CH3NH3PbI3-xClx

1.075 ± 0.01 ns

0.074 ± 0.006

3.3 ± 0.2 ns

ITO / sp-TiO2/ CH3NH3PbI3-xClx

897 ± 10 ps

0.234 ± 0.006

3.32 ± 0.03 ns

ITO / b-TiO2/ CH3NH3PbI3-xClx

715 ± 4 ps

0.121 ± 0.003

3.32 ± 0.05 ns

ITO / m-TiO2/ CH3NH3PbI3-xClx

482 ± 1 ps

A representative PL decay trace for each sample preparation is displayed in Figure 9. The complete table of fitting parameters and figures displaying all traces can be found in the Supporting Information (Figure S6A-E, Table S2). The only TRPL traces that were best described with a single exponential decay were those of CH3NH3PbI3-xClx on m-TiOx. In addition, m-TiOx/CH3NH3PbI3-xClx samples displayed the fastest decay with an average time constant of 482 ± 1 ps. CH3NH3PbI3-xClx films on b-TiO2 and sp-TiO2 displayed slower decay time scales for their first exponential component (715 ± 4 ps and 897 ± 10 ps, respectively). Control films of CH3NH3PbI3-xClx on glass and bare ITO displayed the slowest initial decay in their TRPL traces with time scales slightly above 1 ns. All samples except m-TiOx/CH3NH3PbI3xClx

displayed an additional, longer decay with similar time scales of 3-3.3 ns.


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Figure 9. Normalized TRPL profiles and fitting curves for CH3NH3PbI3-xClx deposited atop the various ETLs. TRPL studies on the CH3NH3PbI3-xClx layer deposited on bare glass and bare ITO were performed as control experiments.

The TRPL results can be better understood by placing them in context with the rest of the experimental data presented in this manuscript. The samples with the largest trap density (m-TiOx/CH3NH3PbI3-xClx) also showed the largest change in their TRPL decay, with the fastest decay time constant (482 ps) and did not display the longer decay component (3-3.3 ns) present in all the other samples.

While not

immediately obvious from the TRPL traces (Figure 9), quantitatively fitting the measured decay dynamics reveals that b-TiO2 and sp-TiO2 ETLs have TRPL decay time scales shorter than those for the bare glass and ITO-coated glass substrates (715 ps and 897 ps vs 1 ns) – although these differences are smaller than the one observed for the samples with a m-TiOx ETL. We interpret the faster fluorescence decay in the CH3NH3PbI3-xClx films atop titania ETLs to reflect the ability of excitons (or photogenerated


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free charges) to undergo charge transfer to the ETLs. Importantly, the different ETLs show substantial variations in their fluorescence dynamics which we attribute to a combination of differences in their trap densities and their energy level alignment with the CH3NH3PbI3-xClx active layer.

CONCLUSIONS We have investigated the morphology, optoelectronic properties, and device characteristics of CH3NH3PbI3-xClx solar cells incorporating titania ETLs fabricated using three different methods: solution processed compact b-TiO2, solution processed mesoscopic m-TiOx, and sputter deposited compact sp-TiO2. By investigating the physical and electronic properties of our titania layers, we were able to systematically elucidate how changes in their morphology, defect concentrations, and energy level alignment affect their performance as effective ETLs in perovskite solar cells. We observe that thin films comprising sp-TiO2 yield fewer defects and a smoother compact film. The fabrication of ETLs having these characteristics is of paramount importance to ensure high power conversion efficiencies when the ETLs are further incorporated into functional devices. As such, we have demonstrated that the efficiency of CH3NH3PbI3-xClx solar cells are mainly limited by trap (defect) and energy level misalignment effects triggered by the ETLs. We do not believe the morphology of the ETL plays an important role on the observed PCEs. From our studies, we have found that sputter-deposited TiO2 is a more efficient electron transport layer for CH3NH3PbI3-xClx solar cells than any of the solution-processed TiO2 layers fabricated in this work. We determined that the superior performance of sputter-deposited TiO2 ETLs is due to its optimized energy level alignment with the CH3NH3PbI3-xClx layer and increased electron transport. These factors, combined with the low trap density in the sputterdeposited TiO2 ETLs, prevent recombination and lead to an increase in the short circuit density, open circuit voltage, and fill factor of the devices thus, leading to higher PCEs.


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EXPERIMENTAL SECTION ETL Fabrication Patterned indium tin oxide (ITO)-coated glass substrates were cleaned in sequentially sonicated baths of soapy water, water, and then isopropyl alcohol. Following sonication, substrates were oxygen plasma cleaned for 30-45 minutes. Solution processed ETLs, b-TiO2 and m-TiO2, were fabricated via a dip-coating process.86-88 Sol-gel precursor solutions were prepared according to previous work.86 Briefly, the sol-gel precursor was prepared by combining a molar ratio of 1 : 16 : 0.4 : 1 titanium(IV)isopropoxide : EtOH : H2O : HCl. The precursor solution was aged for 48 hours at room temperature, and then diluted with ethanol to achieve the desired film thicknesses with dip-coating deposition. For casting 30 nm b-TiO2 thin films, the sol-gel precursor solution was diluted to a 30% by volume solution in ethanol. For m-TiOx films, a 40% by weight Pluronic

®

P123

surfactant,

a

poly(ethyleneoxide)20–poly(propyleneoxide)70–

poly(ethyleneoxide)20 block copolymer, was added to a 20% by volume dilution of the precursor solution. All films were dipped at a speed of 12.6 cm/minute and excess solvent was immediately removed via a hot air stream. The as-cast b-TiO2 films were subsequently annealed in air at 450 oC for 15 mins. To fabricate m-TiOx thin films, a 1:9 dilution of the TiO2 sol-gel precursor to ethanol was used to deposit a 10 nm-thick blocking layer over the ITO-coated glass substrate. A 20 nm-thick layer of m-TiOx was deposited atop the blocking layer. After dipcoating, the as-cast m-TiOx thin films were annealed at 450 oC for 15 mins to yield a porous TiOx layer. The sp-TiO2 thin film was deposited via RF magnetron sputtering deposition using a Denton Discovery 18 sputter system. The sputtering target was a 2” TiO2 disc that is 0.125” thick and it was used as received from Kurt J. Lesker Company (target purity 99.99%). These sputtered films were deposited at room temperature at an applied power of 150 W under a flow


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of Argon at a rate of 2.5 nm/min and a pressure of 4.2 mtorr. No additional gases (i.e., reactive oxygen gas) were introduced in the sputtering chamber during deposition. The sp-TiO2 film thicknesses were controlled by adjusting the sputtering time and deposition rate. The sample thicknesses were monitored by using a quartz crystal microbalance placed close to the ITOcoated glass substrates in the sputtering chamber.

Prior to solar cell fabrication, sp-TiO2

samples were annealed at 450 oC for 4 hours. All procedures were optimized to yield 30nmthick sol-TiO2, m-TiO2, and sp-TiO2 films. Characterization The ETL morphologies were characterized using a Bruker Dimension Icon atomic force microscope (AFM) operated in tapping mode. The perovskite morphologies were examined using a scanning electron microscope (SEM) equipped with a field-emission gun (FEI Quanta 600F) and operated at an accelerating voltage of 10 keV. A Bruker D8 diffractometer with Cu Kα radiation (λ = 1.5418 Å) was used to determine the crystal structure and phase purity of the titania and perovskite layers. The operating voltage and current were kept at 40 kV and 40 mA, respectively. Further morphology and structure characterizations were performed via grazing incidence wide angle X-ray scattering (GIWAXS) studies. These GIWAXS experiments were conducted at the Cornell High Energy Synchrotron Source (CHESS, beamline D1). The x-ray beam energy of 10.6 keV was selected with synthetic multilayer optics (Mo/B4C, 30 Å dspacing). The X-ray beam was aligned above the film’s critical angle and below that of the substrate’s, at 0.25° with respect to the substrate. The scattered intensity was collected with a two-dimensional CCD detector comprising 1024 x 1024 pixels with a size of 46.9 microns. Sample to detector distance was 173 mm.53, 89 All GIWAXS images were background subtracted, and polarization and absorption corrections were applied, though these corrections were


33


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generally small. Additionally, GIXSGUI was used as the graphical user interface to visualize and process the GIWAXS data.90 The optical bandgaps of the TiO2 thin films were determined via UV-Vis-NIR absorption spectroscopy using a Hitachi U4100 UV-Vis-Near IR spectrometer. Qualitative determination of defects within the TiO2 thin films was done via fluorescence spectroscopy (Hitachi F-7000 Fluorescence spectrophotometer) in emission mode. For this ETL defect analysis, an excitation energy of 3.9 eV (320 nm) was used. Here, ETL layers were prepared on Si wafers to avoid interaction with the glass substrate. Reflectance vs. photon energy spectra were used to estimate the absorbance onsets as well as optical bandgaps. Here, the reﬂectance data generated were converted to absorption using a Kubelka−Munk [F(R)hν]n function.(91-93) The intercept of the tangent of the inflexion point of [F(R)hν]n vs photon energy with the x-axis on the Tauc plots was used to estimate the optical bandgaps. Here, n = 1/2 was used to determine the bandgap for a TiO2 having an indirect allowed interband transition. Concentration of charged defects in TiO2 thin films was determined via capacitance-voltage (C-V) studies. For C-V measurements, ETLs were deposited on patterned ITO-coated glass substrates to ensure a device area equivalent to that of the solar cell devices (0.18 cm2).

Silver electrodes were deposited using thermal

evaporation. A 0.45 mm2 Hg probe was used to contact the silver electrode within the C-V device. A Keithley 4200 semiconductor characterization system was used to measure the capacitance of the devices as a function of voltage. To compare the magnitude of the midgap voltage difference (∆Vmg), we normalized the C-V characteristics with respect to the sample the largest hysteresis. For ultraviolet photoelectron spectroscopy (UPS) measurements, all films fabricated were handled in inert N2 atmosphere and never exposed to oxygen and water levels above 5 ppm before introduction into the ultra-high vacuum (UHV) system. UPS spectra were


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acquired in a custom-built UHV chamber equipped with a cylindrical mirror electron analyzer and operated at a base pressure of 10-10 Torr. UPS was performed by using He I (21.22 eV) and He II (40.8 eV) excitation lines of a He plasma discharge lamp. Spectra were taken at a pass energy of 5 eV for a nominal experimental resolution smaller than 150 meV. The calibration of the binding energy scale of our UPS setups was performed by acquiring an UPS spectra for sputter cleaned Au and Ag thin films. Such calibration yielded binding energy values for Ag 3d5/2 and Au 4f7/2 peaks at about 368.27 eV and 84.00 eV, respectively. These numbers agree with the standard tabulated values defined by the International Organization of Standardization 15472 to within ± 0.02 eV. As a clear advantage, the Fermi level of a metal serves as an ideal choice as it can be determined using UPS and also suitably functions as the zero of the binding energy scale. Thus, all energy levels acquired in our studies are referenced to a common Fermi Level (0 eV) which is related to the Fermi level of a known metal (either Au or Ag). All measurements were carried out at normal take-off angles. Time-resolved photoluminescence (TRPL) studies were conducted using a time-correlated single-photon counting (TCSPC) fluorescence setup. A Becker & Hickl 405 nm picosecond diode laser was used as an excitation source, with a 405 nm bandpass excitation filter. Crossed polarizers and reflective neutral density filters were used to control the power of the excitation beam. For these experiments, the repetition rate of the laser was set to 20 MHz, and the beam at the sample position had a power of approximately 100 µW with a spot size of 0.25 cm2. Fluorescence was collected and collimated using spherical lens in a direction perpendicular to the excitation. A 500 nm longpass filter was used to reject scattered pump photons, and the fluorescence signal was focused onto a cooled photomultiplier tube (Hamamatsu, H7422P-40) for detection. Time-correlated detection was performed with a Becker & Hickl SPC-630 card.


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Samples were deposited on glass slides, with or without ITO and encapsulated using UV-curing Norland Optical Adhesive 61 to avoid degradation. Samples were placed at a 45 degree angle with respect to the excitation beam to reduce the detection of reflected pump photons, and the excitation was directed onto the sample through the glass slide side of the film stack. The setup had an instrument response function (IRF) with a full-width at half-maximum of 420 ps. The IRF was convoluted with a single- or bi-exponential function in a nonlinear least-squares fitting algorithm (Matlab) to obtain the decay time constants and amplitudes for each sample. The uncertainties reported in the fitting parameters are the result of error propagation through this nonlinear fitting algorithm (one standard deviation). Solar cell fabrication ITO-coated glass substrates and ETLs were clean and prepared as stated above. To avoid oxygen and moisture, the substrates were transferred to a glovebox.

Methylammonium lead

iodide/chloride (CH3NH3PbI3-xClx) solutions were prepared with a 3:1 ratio of methylammonium iodide:PbCl2 in dimethylformamide. The concentration of Pb was kept to 0.8 M. Perovskite solutions were spun-cast atop the ETL coated ITO glass substrate under inert atmosphere at 3000 rpm for 40 s. The perovskite layer formed during spincoating was annealed at 100 °C for 2 hrs. Afterwards, the hole transporting layer (HTL) was deposited by spincoating at 2000 rpm for 120s.

The HTL recipe was prepared by dissolving 10 mg of poly[bis(4-phenyl)(2,4,6-

trimethylphenyl)amine] (PTAA), 7.5 µL of 4-tert-butylpyridine (TBP), and 7.5 µL of a lithium bis-(trifluromethylsulfonyl)imide (Li-TFSI) solution (170 mg of Li-TFSI in 1 mL of acetonitrile). This mixture was dissolved in 1 mL toluene. To complete the solar cells, a 100-nm thick layer of Au was deposited by thermal evaporation through stencil masks as top electrodes. The active area for each device was 0.18 cm2. All device measurements, including IV


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characteristics and PCE measurements, external quantum efficiency (EQE) and electrical impedance spectroscopy (EIS), were conducted in air. IV characteristics and EQE measurements were taken on a Solar Cell Quantum Efficiency Measurement System (PV measurements, model QEXL). PCE measurements were taken under AM 1.5G 100 mW/cm2 illumination. IV characteristics were measured by running a forward scan (1 to -0.2 V) followed by a reverse scan (-0.2 to 1 V) at a scan rate of 0.37 V/s. For EQE measurements, samples were illuminated with monochromatic light in 10 nm increments from 300 to 1100 nm. AC EIS measurements were taken on a CH Instruments Electrochemical Workstation. A bias equivalent to the Voc of each device was applied during measurement. Frequency ranges from 0.1 MHz to 200-500 Hz were used. EIS spectra were fit to the equivalent electrical circuit in Figure 8 using a CHI Version 14.09 software. Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI:xx.xxxx/acsami.xxxxxx. Details on the structural and electronic characterization of the ETLs; XRD spectra for b-TiO2, mTiOx, and sp-TiO2 ETLs; UPS spectra for pristine b-TiO2, m-TiOx, and sp-TiO2 ETLs on ITO as well as the UPS spectra for a 500nm-thick CH3NH3PbI3 layer atop the various ETLs; device statistics and J-V curves for CH3NH3PbI3-xClx solar cells atop b-TiO2, m-TiOx, and sp-TiO2; dark curves for CH3NH3PbI3-xClx solar cells deposited atop b-TiO2, m-TiOx, and sp-TiO2 ETLs; TRPL profiles and fitting curves for CH3NH3PbI3-xClx thin films atop bare glass, ITO-coated glass substrates, b-TiO2, m-TiOx, and sp-TiO2 ETLs. AUTHOR INFORMATION


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Corresponding author *Email: [email protected] Notes The authors declare no competing financial interest. Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ACKNOWLEDGMENT This work was supported by the NSF MRSEC program at the University of Utah under grant number DMR-1121252. GIWAXS studies were conducted at CHESS, which is supported by NSF and NIH/NIGMS under NSF award DMR-1332208. WN acknowledges funding from the Utah Governor’s Office of Energy Development. LWB would also like to acknowledge the financial support from the Marion Milligan Mason Award for Women in the chemical sciences administered by the American Association for the Advancement of Science (AAAS).

REFERENCES (1)

Zhou, H.; Chen, Q.; Li, G.; Luo, S.; Song, T.-b.; Duan, H.-s.; Hong, Z.; You, J.; Liu, Y.; Yang, Y. Interface Engineering of Highly Efficient Perovskite Solar Cells. Science 2014, 345, 542-546. (2) Nature Materials Editorial, Nat. Mater. 2014, 13, 837-837. (3) Gratzel, M. The Light and Shade of Perovskite Solar Cells. Nat. Mater. 2014, 13, 838842. (4) McGehee, M. Perovskite Solar Cells: Continuing to Soar. Nat. Mater. 2014, 13, 845-846. (5) National Renewable Energy Laboratory. Efficiency_Chart http://Www.Nrel.Gov/Ncpv/Images/Efficiency_Chart.Jpg Accessed August 2017. (6) Mitzi, D. in Progress in Inorganic Chemistry, John Wiley & Sons, Inc., 2007, pp. 1-121.


38

Page 39 of 45 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


(7)

(8) (9)

(10)

(11)

(12)

(13)

(14)

(15)

(16)

(17)

(18)

(19) (20)

Baikie, T.; Fang, Y.; Kadro, J. M.; Schreyer, M.; Wei, F.; Mhaisalkar, S. G.; Graetzel, M.; White, T. J. Synthesis and Crystal Chemistry of the Hybrid Perovskite (CH3NH3)PbI3 for Solid-State Sensitised Solar Cell Applications. J. Mater. Chem. A 2013, 1, 5628-5641. Umari, P.; Mosconi, E.; De Angelis, F. Relativistic GW Calculations on CH3NH3PbI3 and CH3NH3SnI3 Perovskites for Solar Cell Applications. Sci. Rep. 2014, 4, 4467-4474. Tanaka, K.; Takahashi, T.; Ban, T.; Kondo, T.; Uchida, K.; Miura, N. Comparative Study on the Excitons in Lead-Halide-Based Perovskite-Type Crystals CH3NH3PbBr3 CH3NH3PbI3. Solid State Commun. 2003, 127, 619-623. Kulkarni, S. A.; Baikie, T.; Boix, P. P.; Yantara, N.; Mathews, N.; Mhaisalkar, S. BandGap Tuning of Lead Halide Perovskites Using a Sequential Deposition Process. J. Mater. Chem. A 2014, 2, 9221-9225. Ponseca, C. S.; Savenije, T. J.; Abdellah, M.; Zheng, K.; Yartsev, A.; Pascher, T.; Harlang, T.; Chabera, P.; Pullerits, T.; Stepanov, A.; Wolf, J.-P.; Sundström, V. Organometal Halide Perovskite Solar Cell Materials Rationalized: Ultrafast Charge Generation, High and Microsecond-Long Balanced Mobilities, and Slow Recombination. J. Am. Chem. Soc. 2014, 136, 5189-5192. Stranks, S. D.; Eperon, G. E.; Grancini, G.; Menelaou, C.; Alcocer, M. J. P.; Leijtens, T.; Herz, L. M.; Petrozza, A.; Snaith, H. J. Electron-Hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber. Science 2013, 342, 341344. Xing, G.; Mathews, N.; Sun, S.; Lim, S. S.; Lam, Y. M.; Grätzel, M.; Mhaisalkar, S.; Sum, T. C. Long-Range Balanced Electron- and Hole-Transport Lengths in OrganicInorganic CH3NH3PbI3. Science 2013, 342, 344-347. Wehrenfennig, C.; Liu, M.; Snaith, H. J.; Johnston, M. B.; Herz, L. M.; Charge-Carrier Dynamics in Vapour-Deposited Films of the Organolead Halide Perovskite CH3NH3PbI3−xClx. Energy Environ. Sci. 2014, 7, 2269-2275. Wehrenfennig, C.; Eperon, G. E.; Johnston, M. B.; Snaith, H. J.; Herz, L. M.; High Charge Carrier Mobilities and Lifetimes in Organolead Trihalide Perovskites. Adv. Mater. 2014, 26, 1584-1589. Chen, S.; Lei, L.; Yang, S.; Liu, Y.; Wang, Z.-S. Characterization of Perovskite Obtained from Two-Step Deposition on Mesoporous Titania. ACS Appl. Mater. Interfaces 2015, 7, 25770-25776. Hill, A. H.; Smyser, K. E.; Kennedy, C. L.; Massaro, E. S.; Grumstrup, E. M. Screened Charge Carrier Transport in Methylammonium Lead Iodide Perovskite Thin Films. J. Phys. Chem. Lett. 2017, 8, 948-953. Shi, D.; Adinolfi, V.; Comin, R.; Yuan, M.; Alarousu, E.; Buin, A.; Chen, Y.; Hoogland, S.; Rothenberger, A.; Katsiev, K.; Losovyj, Y.; Zhang, X.; Dowben, P. A.; Mohammed, O. F.; Sargent, E. H.; Bakr, O. M. Low Trap-State Density and Long Carrier Diffusion in Organolead Trihalide Perovskite Single Crystals. Science 2015, 347, 519-522. Liu, M.; Johnston, M. B.; Snaith, H. J. Efficient Planar Heterojunction Perovskite Solar Cells by Vapour Deposition. Nature 2013, 501, 395-398. Liu, X.; Yu, H.; Yan, L.; Dong, Q.; Wan, Q.; Zhou, Y.; Song, B.; Li, Y. Triple Cathode Buffer Layers Composed of PCBM, C60, and LiF for High-Performance Planar Perovskite Solar Cells. ACS Appl. Mater. Interfaces 2015, 7, 6230-6237.


39


(21)

(22) (23)

(24)

(25)

(26)

(27)

(28)

(29)

(30)

(31)

(32)

(33)

(34)

Page 40 of 45

Wang, Q.; Shao, Y.; Dong, Q.; Xiao, Z.; Yuan, Y.; Huang, J. Large Fill-Factor Bilayer Iodine Perovskite Solar Cells Fabricated by a Low-Temperature Solution-Process. Energy Environ. Sci. 2014, 7, 2359-2365. Yang, G.; Tao, H,; Qin, P.; Ke, W.; Fang, G. Recent Progress in Electron Transport Layers for Efficient Perovskite Solar Cells. J. Mater. Chem. A 2016, 4, 3970-3990. Correa-Baena, J.-P.; Abate, A.; Saliba, M.; Tress, W.; Jesper Jacobsson, T.; Gratzel, M.; Hagfeldt, A. The Rapid Evolution of Highly Efficient Perovskite Solar Cells. Energy Environ. Sci. 2017, 10, 710-727. Leijtens, T.; Eperon, G. E.; Pathak, S.; Abate, A.; Lee, M. M.; Snaith, H. J. Overcoming Ultraviolet Light Instability of Sensitized TiO₂ with Meso-Superstructured Organometal Tri-Halide Perovskite Solar Cells. Nat. Commun. 2013, 4, 2885-2893. Ito, S.; Tanaka, S.; Manabe, K.; Nishino, H. Effects of Surface Blocking Layer of Sb2S3 on Nanocrystalline TiO2 for CH3NH3PbI3 Perovskite Solar Cells. J. Phys. Chem. C 2014, 118, 16995-17000. Snaith, H. J.; Abate, A.; Ball, J. M.; Eperon, G. E.; Leijtens, T.; Noel, N. K.; Stranks, S. D.; Wang, J. T.-W.; Wojciechowski, K.; Zhang, W. Anomalous Hysteresis in Perovskite Solar Cells. J. Phys. Chem. Lett. 2014, 5, 1511-1515. Giordano, F.; Abate, A.; Correa Baena, J. P.; Saliba, M.; Matsui, T.; Im, S. H.; Zakeeruddin, S. M.; Nazeeruddin, M. K.; Hagfeldt, A.; Graetzel, M. Enhanced Electronic Properties in Mesoporous TiO2 via Lithium Doping for High-Efficiency Perovskite Solar Cells. Nat. Commun. 2016, 7, 10379-10385. Yang, M.; Guo, R.; Kadel, K.; Liu, Y.; O'Shea, K.; Bone, R.; Wang, X.; He, J.; Li, W. Improved Charge Transport of Nb-doped TiO2 Nanorods in Methylammonium Lead Iodide Bromide Perovskite Solar Cells. J. Mater. Chem.A 2014, 2, 19616-19622. Wang, H.-H.; Chen, Q.; Zhou, H.; Song, L.; Louis, Z. S.; Marco, N. D.; Fang, Y.; Sun, P.; Song, T.-B.; Chen, H.; Yang, Y. Improving the TiO2 Electron Transport Layer in Perovskite Solar Cells Using Acetylacetonate-Based Additives. J. Mater. Chem. A 2015, 3, 9108-9115. Chandiran, A.K.; Abdi-Jalebi, M.; Nazeeruddin, M. K.; Grätzel, M. Analysis of Electron Transfer Properties of ZnO and TiO2 Photoanodes for Dye-Sensitized Solar Cells. ACS Nano 2014, 8, 2261-2268. Dong, X.; Hu, H.; Lin, B.; Ding, J.; Yuan, N. The Effect of ALD-ZnO Layers on the Formation of CH3NH3PbI3 with Different Perovskite Precursors and Sintering Temperatures. Chem. Commun. 2014, 50, 14405-14408. Ke, W.; Fang, G.; Liu, Q.; Xiong, L.; Qin, P.; Tao, H.; Wang, J.; Lei, H.; Li, B.; Wan, J.; Yang, G.; Yan, Y. Low-Temperature Solution-Processed Tin Oxide as an Alternative Electron Transporting Layer for Efficient Perovskite Solar Cells. J. Am. Chem. Soc. 2015, 137, 6730-6733. Conings, B.; Baeten, L.; Jacobs, T.; Dera, R.; D’Haen, J.; Manca, J.; Boyen, H. G. An Easy-to-Fabricate Low-Temperature TiO2 Electron Collection Layer for High Efficiency Planar Heterojunction Perovskite Solar Cells. APL Mater. 2014, 2, 081505. Wojciechowski, K.; Saliba, M.; Leijtens, T.; Abate, A.; Snaith, H. J. Sub-150 °C Processed Meso-Superstructured Perovskite Solar Cells with Enhanced Efficiency. Energy Environ. Sci. 2014, 7, 1142-1147.


40

Page 41 of 45 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


(35)

(36)

(37)

(38)

(39)

(40)

(41)

(42)

(43)

(44)

(45)

(46)

(47)

(48)

Dharani, S.; Mulmudi, H. K.; Yantara, N.; Thu Trang, P. T.; Park, N. G.; Graetzel, M.; Mhaisalkar, S.; Mathews, N.; Boix, P. P. High Efficiency Electrospun TiO2 Nanofiber Based Hybrid Organic–Inorganic Perovskite Solar Cell. Nanoscale 2014, 6, 1675-1679. Heo, J. H.; Im, S. H.; Noh, J. H.; Mandal, T. N.; Lim, C.-S.; Chang, J. A.; Lee, Y. H.; Kim, H.-J.; Sarkar, A.; NazeeruddinMd, K.; Gratzel, M.; Seok, S. I. Efficient Inorganic– Organic Hybrid Heterojunction Solar Cells Containing Perovskite Compound and Polymeric Hole Conductors. Nat. Photonics 2013, 7, 486-491. Zhao, Y.; Zhu, K. CH3NH3Cl-Assisted One-Step Solution Growth of CH3NH3PbI3: Structure, Charge-Carrier Dynamics, and Photovoltaic Properties of Perovskite Solar Cells. J. Phys. Chem. C 2014, 118, 9412-9418. Chen, C.; Cheng, Y.; Dai, Q.; Song, H. Radio Frequency Magnetron Sputtering Deposition of TiO2 Thin Films and Their Perovskite Solar Cell Applications. Sci. Rep. 2015, 5, 17684-17696. Ke, W.; Fang, G.; Wang, J.; Qin, P.; Tao, H.; Lei, H.; Liu, Q.; Dai, X.; Zhao, X. Perovskite Solar Cell with an Efficient TiO2 Compact Film. ACS Appl. Mater. Interfaces 2014, 6, 15959-15965. Mali, S. S.; Hong, C. K.; Inamdar, A. I.; Im, H.; Shim, S. E. Efficient Planar n-i-p Type Heterojunction Flexible Perovskite Solar Cells with Sputtered TiO2 Electron Transporting Layers. Nanoscale 2017, 9, 3095-3104. Hong, S.; Han, A.; Lee, E. C.; Ko, K.-W.; Park, J.-H.; Song, H.-J.; Han, M.-H.; Han, C.H. A Facile and Low-Cost Fabrication of TiO2 Compact Layer for Efficient Perovskite Solar Cells. Curr. Appl. Phys. 2015, 15, 574-579. Chandiran, A. K.; Yella, A.; Mayer, M. T.; Gao, P.; Nazeeruddin, M. K.; Grätzel, M. Sub-Nanometer Conformal TiO2 Blocking Layer for High Efficiency Solid-State Perovskite Absorber Solar Cells. Adv. Mater. 2014, 26, 4309-4312. Hu, H.; Dong, B.; Hu, H.; Chen, F.; Kong, M.; Zhang, Q.; Luo, T.; Zhao, L.; Guo, Z.; Li, J.; Xu, Z.; Wang, S.; Eder, D.; Wan, L. Atomic Layer Deposition of TiO2 for a HighEfficiency Hole-Blocking Layer in Hole-Conductor-Free Perovskite Solar Cells Processed in Ambient Air. ACS Appl. Mater. Interfaces 2016, 8, 17999-18007. Wu, Y.; Yang, X.; Chen, H.; Zhang, K.; Qin, C.; Liu, J.; Peng, W.; Islam, A.; Bi, E.; Ye, F. Highly Compact TiO2 Layer for Efficient Hole-Blocking in Perovskite Solar Cells. Appl. Phys. Express 2014, 7, 052301. Kim, I. S.; Haasch, R. T.; Cao, D. H.; Farha, O. K.; Hupp, J. T.; Kanatzidis, M. G.; Martinson, A. B. F. Amorphous TiO2 Compact Layers via ALD for Planar Halide Perovskite Photovoltaics. ACS Appl. Mater. Interfaces 2016, 8, 24310-24314. Chung, C.-C.; Lee, C. S.; Jokar, E.; Kim, J. H.; Diau, E. W.-G., Well-organized mesoporous tio2 photoanode by using amphiphilic graft copolymer for efficient perovskite solar cells. J. Phys. Chem. C 2016, 120, 9619-9627. Whittaker-Brooks, L.; Mativetsky, J. M.; Woll, A.; Smilgies, D.; Loo, Y.-L. Sputtered ZnO Seed Layer Enhances Photovoltaic Behavior in Hybrid ZnO/P3HT Solar Cells. Org. Electron. 2013, 14, 3477-3483. Li, D.; Ohashi, N.; Hishita, S.; Kolodiazhnyi, T.; Haneda, H. Origin of Visible-LightDriven Photocatalysis: A Comparative Study on N/F-Doped and N–F-Codoped TiO2 Powders by Means of Experimental Characterizations and Theoretical Calculations. J. Solid State Chem. 2005, 178, 3293-3302.


41


(49)

(50)

(51)

(52)

(53)

(54)

(55)

(56)

(57)

(58) (59)

(60)

(61) (62)

(63)

Page 42 of 45

Lei, Y.; Zhang, L. D.; Meng, G. W.; Li, G. H.; Zhang, X. Y.; Liang, C. H.; Chen, W.; Wang, S. X.; Preparation and Photoluminescence of Highly Ordered TiO2 Nanowire Arrays. Appl. Phys. Lett. 2001, 78, 1125-1127. Bok Kim, J.; Ahn, S.; Ju Kang, S.; Nuckolls, C.; Loo, Y.-L. Ligand Chemistry of Titania Precursor Affects Transient Photovoltaic Behavior in Inverted Organic Solar Cells. Appl. Phys. Lett. 2013, 102, 103302. Shao, J.; Yang, S.; Lei, L.; Cao, Q.; Yu, Y.; Liu, Y., Pore size dependent hysteresis elimination in perovskite solar cells based on highly porous tio2 films with widely tunable pores of 15–34 nm. Chem. Mater. 2016, 28, 7134-7144. Tan, K. W.; Moore, D. T.; Saliba, M.; Sai, H.; Estroff, L. A.; Hanrath, T.; Snaith, H. J.; Wiesner, U. Thermally Induced Structural Evolution and Performance of Mesoporous Block Copolymer-Directed Alumina Perovskite Solar Cells. ACS Nano 2014, 8, 47304739. Whittaker-Brooks, L.; McClain, W. E.; Schwartz, J.; Loo, Y.-L. Donor-Acceptor Interfacial Interactions Dominate Device Performance in Hybrid P3HT-ZnO NanowireArray Solar Cells. Adv. Energy Mater.2014, 4, 1400585. Guerrero, A.; Juarez-Perez, E. J.; Bisquert, J.; Mora-Sero, I.; Garcia-Belmonte, G. Electrical Field Profile and Doping in Planar Lead Halide Perovskite Solar Cells. Appl. Phys. Lett. 2014, 105, 133902. Tress, W.; Marinova, N.; Moehl, T.; Zakeeruddin, S. M.; Nazeeruddin, M. K.; Gratzel, M. Understanding the Rate-Dependent J–V Hysteresis, Slow Time Component, and Aging in CH3NH3PbI3 Perovskite Solar Cells: the Role of a Compensated Electric Field. Energy Environ. Sci. 2015, 8, 995-1004. Schulz, P.; Edri, E.; Kirmayer, S.; Hodes, G.; Cahen, D.; Kahn, A. Interface Energetics in Organo-Metal Halide Perovskite-Based Photovoltaic Cells. Energy Environ. Sci. 2014, 7, 1377-1381. Schulz, P.; Whittaker-Brooks, L. L.; MacLeod, B. A.; Olson, D. C.; Loo, Y.-L.; Kahn, A. Electronic Level Alignment in Inverted Organometal Perovskite Solar Cells. Adv. Mater. Interfaces 2015, 2, 1400532. Shao, Y.; Yuan, Y.; Huang, J., Correlation of energy disorder and open-circuit voltage in hybrid perovskite solar cells. Nature Energy 2016, 1, 15001. Zeng, W.; Liu, X.; Guo, X.; Niu, Q.; Yi, J.; Xia, R.; Min, Y., Morphology analysis and optimization: Crucial factor determining the performance of perovskite solar cells. Molecules 2017, 22. Lim, K.-G.; Ahn, S.; Kim, Y.-H.; Qi, Y.; Lee, T.-W., Universal energy level tailoring of self-organized hole extraction layers in organic solar cells and organic-inorganic hybrid perovskite solar cells. Energy Environ. Sci. 2016, 9, 932-939. Ou, Q.-D.; Li, C.; Wang, Q.-K.; Li, Y.-Q.; Tang, J.-X. Recent Advances in Energetics of Metal Halide Perovskite Interfaces. Adv. Mater. Interfaces 2017, 4, 1600694. Dualeh, A.; Moehl, T.; Tétreault, N.; Teuscher, J.; Gao, P.; Nazeeruddin, M. K.; Grätzel, M. Impedance Spectroscopic Analysis of Lead Iodide Perovskite-Sensitized Solid-State Solar Cells. ACS Nano 2013, 8, 362-373. Unger, E. L.; Hoke, E. T.; Bailie, C. D.; Nguyen, W. H.; Bowring, A. R.; Heumuller, T. Christoforo, M. G.; McGehee, M. D. Hysteresis and Transient Behavior in Current– Voltage Measurements of Hybrid-Perovskite Absorber Solar Cells. Energy Environ. Sci. 2014, 7, 3690-3698.


42

Page 43 of 45 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


(64) (65)

(66)

(67)

(68)

(69)

(70)

(71)

(72) (73)

(74)

(75) (76)

(77) (78)

(79)

Frost, J. M.; Butler, K. T.; Walsh, A. Molecular Ferroelectric Contributions to Anomalous Hysteresis in Hybrid Perovskite Solar Cells. APL Mater. 2014, 2, 081506. Shao, Y.; Xiao, Z.; Bi, C.; Yuan, Y.; Huang, J. Origin and Elimination of Photocurrent Hysteresis by Fullerene Passivation in CH3NH3PbI3 Planar Heterojunction Solar Cells. Nat. Commun. 2014, 5, 7081. Kim, H.-S.; Park, N.-G. Parameters Affecting I–V Hysteresis of CH3NH3PbI3 Perovskite Solar Cells: Effects of Perovskite Crystal Size and Mesoporous TiO2 Layer. J. Phys. Chem. Lett. 2014, 5, 2927-2934. Sanchez, R. S.; Gonzalez-Pedro, V.; Lee, J.-V.; Park, N.-G.; Kang, Y. S.; Mora-Sero, I.; Bisquert, J. Slow Dynamic Processes in Lead Halide Perovskite Solar Cells. Characteristic Times and Hysteresis. J. Phys. Chem. Lett. 2014, 5, 2357-2363. Chen, B.; Yang, M.; Zheng, X.; Wu, C.; Li, W.; Yan, Y.; Bisquert, J.; Garcia-Belmonte, G.; Zhu, K.; Priya, S. Impact of Capacitive Effect and Ion Migration on the Hysteretic Behavior of Perovskite Solar Cells. J. Phys. Chem. Lett. 2015, 6, 4693-4700. Frost, J. M.; Butler, K. T.; Brivio, F.; Hendon, C. H.; van Schilfgaarde, M.; Walsh, A. Atomistic Origins of High-Performance in Hybrid Halide Perovskite Solar Cells. Nano Lett. 2014, 14, 2584-2590. Eames, C.; Frost, J. M.; Barnes, P. R. F.; O'Regan, B. C.; Walsh, A.; Islam, M. S. Ionic Transport in Hybrid Lead Iodide Perovskite Solar Cells. Nat. Commun. 2015, 6, 74977505. Meloni, S.; Moehl, T.; Tress, W.; Franckevičius, M.; Saliba, M.; Lee,Y. H.; Gao, P.; Nazeeruddin, M. K.; Zakeeruddin, S. M.; Rothlisberger, U.; Graetzel, M. Ionic Polarization-Induced Current–Voltage Hysteresis in CH3NH3PbX3 Perovskite Solar Cells. Nat. Commun. 2016, 7, 10334-10343. Heiser, T.; Weber, E. R. Transient Ion-Drift-Induced Capacitance Signals in Semiconductors. Phys. Rev. B 1998, 58, 3893-3903. Stoumpos, C. C.; Malliakas, C. D.; Kanatzidis, M. G. Semiconducting Tin and Lead Iodide Perovskites with Organic Cations: Phase Transitions, High Mobilities, and NearInfrared Photoluminescent Properties. Inorg. Chem. 2013, 52, 9019-9038. Brivio, F.; Walker, A. B.; Walsh, A. Structural and Electronic Properties of Hybrid Perovskites for High-Efficiency Thin-Film Photovoltaics from First-Principles. APL Mater. 2013, 1, 042111. Chen, H.-W.; Sakai, N.; Ikegami, M.; Miyasaka, T. Emergence of Hysteresis and Transient Ferroelectric Response in Organo-Lead Halide Perovskite Solar Cells. J. Phys. Wei, J.; Zhao, Y.; Li, H.; Li, G.; Pan, J.; Xu, D.; Zhao, Q.; Yu, D. Hysteresis Analysis Based on the Ferroelectric Effect in Hybrid Perovskite Solar Cells. J. Phys. Chem. Lett. 2014, 5, 3937-3945. Mihailetchi, V. D.; Wildeman, J.; Blom, P. W. M. Space-Charge Limited Photocurrent. Phys. Rev. Lett. 2005, 94, 126602. Koster, L. J. A.; Mihailetchi, V. D.; Xie, H.; Blom, P. W. M. Origin of the Light Intensity Dependence of the Short-Circuit Current of Polymer/Fullerene Solar Cells. Appl. Phys. Lett. 2005, 87, 203502. Bi, D.; Yang, L.; Boschloo, G.; Hagfeldt, A.; Johansson, E. M. J. Effect of Different Hole Transport Materials on Recombination in CH3NH3PbI3 Perovskite-Sensitized Mesoscopic Solar Cells. J. Phys. Chem. Lett. 2013, 4, 1532-1536.


43


(80)

(81)

(82)

(83) (84)

(85)

(86) (87)

(88) (89) (90)

(91)

(92) (93)

Page 44 of 45

Kuik, M.; Koster, L. J. A.; Wetzelaer, G. A. H.; Blom, P. W. M. Trap-Assisted Recombination in Disordered Organic Semiconductors. Phys. Rev. Lett. 2011, 107, 256805. Yang, D.; Zhou, X.; Yang, R.; Yang, Z.; Yu, W.; Wang, X.; Li, C.; Liu, S.; Chang, R. P. H. Surface Optimization to Eliminate Hysteresis for Record Efficiency Planar Perovskite Solar Cells. Energy Environ. Sci. 2016, 9, 3071-3078. Wetzelaer, G.-J. A. H.; Scheepers, M.; Sempere, A. M.; Momblona, C.; Ávila, J.; Bolink, H. J. Trap-Assisted Non-Radiative Recombination in Organic–Inorganic Perovskite Solar Cells. Adv. Mater. 2015, 27, 1837-1841. Lin, Q.; Nagiri, R. C. R.; Burn, P. L.; Meredith, P. Considerations for Upscaling of Organohalide Perovskite Solar Cells. Adv. Opt. Mater. 2017, 5, 201600819. Kim, H.-S.; Mora-Sero, I.; Gonzalez-Pedro, V.; Fabregat-Santiago, F.; Juarez-Perez, E. J.; Park, N.-G.; Bisquert, J. Mechanism of Carrier Accumulation in Perovskite ThinAbsorber Solar Cells. Nat. Commun. 2013, 4, 2242-2249. Kim, H.-S.; Lee, J.-W.; Yantara, N.; Boix, P. P.; Kulkarni, S. A.; Mhaisalkar, S.; Grätzel, M.; Park, N.-G. High Efficiency Solid-State Sensitized Solar Cell-Based on Submicrometer Rutile TiO2 Nanorod and CH3NH3PbI3 Perovskite Sensitizer. Nano Lett. 2013, 13, 2412-2417. Nimens, W. J.; Whittaker-Brooks, L.; Bartl, M. H. Enhanced Sensing in Mixed Porous– Solid Photonic Stacks. J. Mater. Chem. C 2016, 4, 668-672. Barhoum, M.; Morrill, J. M.; Riassetto, D.; Bartl, M. H. Rapid Sol–Gel Fabrication of High-Quality Thin-Film Stacks on Planar and Curved Substrates. Chem. Mater. 2011, 23, 5177-5184. Dahlby, M. R.; Barhoum, M.; Bartl, M. H. Sol–Gel-Derived Thin-Film Stacks with High Radiation Stability. Thin Solid Films 2014, 562, 435-439. L. Whittaker-Brooks, J. Gao, A. K. Hailey, C. R. Thomas, N. Yao, Y.-L. Loo, J. Mater. Chem. C 2015, 3, 2686-2692. Jiang, Z. GIXSGUI: a MATLAB Toolbox for Grazing-Incidence X-ray Scattering Data Visualization and Reduction, and Indexing of Buried Three-Dimensional Periodic Nanostructured Films. J. Appl.Crystallogr. 2015, 48, 917-926. Rahman, A. A.; Huang, R.; Whittaker-Brooks, L., Distinctive Extrinsic Atom Effects on the Structural, Optical, and Electronic Properties of Bi2S3-xSex Solid Solutions. Chem. Mater. 2016, 28, 6544-6552. Nobbs, J. H. Kubelka—Munk Theory and the Prediction of Reflectance. Rev. Prog. Color. Relat. Top., 1985, 15, 66−75. Murphy, A. B. Modified Kubelka−Munk Model for Calculation of the Reflectance of Coatings with Optically-Rough Durfaces. J. Phys. D: Appl. Phys. 2006, 39, 3571.

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Morphology and Optoelectronic Variations Underlying the Nature of

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