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Impact of Titanium Dioxide Surface Defects on the Interfacial Composition and Energetics of Evaporated Perovskite Active Layers R. Clayton Shallcross, Selina Olthof, Klaus Meerholz, and Neal R. Armstrong ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b09935 • Publication Date (Web): 07 Aug 2019 Downloaded from pubs.acs.org on August 8, 2019

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ACS Applied Materials & Interfaces

Impact of Titanium Dioxide Surface Defects on the Interfacial Composition and Energetics of Evaporated Perovskite Active Layers R. Clayton Shallcross,*1 Selina Olthof,2 Klaus Meerholz2 and Neal R. Armstrong1 1

Department of Chemistry and Biochemistry, University of Arizona, Tucson, Arizona 85721, United States of America. Department of Chemistry, University of Cologne, Luxemburger Str. 116, 50939, Cologne, Germany. KEYWORDS: hybrid perovskite, vacuum deposition, stoichiometry, titanium dioxide, surface defects, defect passivation, interface chemistry, energetics 2

ABSTRACT: This investigation elucidates critical Brønsted and Lewis acid-base interactions at the titanium dioxide (TiO2) surface that control the interfacial composition and, thus, energetics of vacuum-processed methylammonium lead iodide (MAPbI3) perovskite active layers (PALs). In situ photoelectron spectroscopy analysis shows that interfacial growth, chemical composition and energetics of co-deposited methylammonium iodide (MAI)/PbI2 thin films is significantly different on bare and (3-aminopropyl)triethoxysilane (APTES)-functionalized TiO2 surfaces. Mass spectroscopy analysis indicates that MAI dissociates into HI and methylamine in the gas-phase and suggests that MAPbI3 nucleation is preceded by adsorption and coupling of these volatile MAI precursors. Prior to MAPbI3 nucleation on the bare TiO2 surface, we suggest that high coverages of methylamine adsorbed to surface defect sites (e.g., undercoordinated Ti atoms and hydroxyls) promote island-like growth of large, PbI2-rich nuclei that inhibit nucleation and lead to a thick substoichiometric interface layer that impedes charge transport and collection energetics. APTES functional groups passivate TiO2 surface defects and facilitate more conformal growth of small, PbI2-rich nuclei that enhance MAPbI3 nucleation and significantly improve interfacial energetics for charge transport and extraction. This work highlights the considerable influence of TiO2 surface chemistry on PAL composition and energetics, which are critical factors that impact the performance and long-stability of these materials in emerging photovoltaic device technologies.

INTRODUCTION Hybrid organic/inorganic perovskite active layers (PALs; e.g., methylammonium lead iodide (MAPbI3) and related formamidinium (FA) and alkali cation lead/tin halides) have recently shown extraordinary performance in thin film photovoltaic (PV) devices, enabling power conversion efficiencies (PCE) in excess of 20% in platforms that appear to be easily scaled to large areas.1-3 PALs are also of interest in emerging photodetector, light-emitting diode, and laser technologies due to their low material cost, ease of processing, long diffusion lengths and high mobilities for charge carriers, strong optical absorption, and compositionally-tuned absorption/luminescence.4 Broader applications of PALs are impeded by several knowledge gaps.5 In particular, recent reports have indicated that the metal oxide (MOx)/PAL interface plays a key role in device performance and degradation.3, 6 X-ray photoelectron spectroscopy (XPS) investigations have shown that substoichiometric interface layers (SILs) precede nucleation of stoichiometric MAPbI3 PALs.7-8 The stoichiometry of MAPbI3 films has been shown to systematically impact the electronic structure;9 therefore, controlling SIL composition is a critical aspect of charge transport and extraction/injection processes at the PAL/MOx interface. Migration of amorphous material associated with SILs on prototypical titanium dioxide (TiO2) contacts has also been linked to current-voltage (J-V) hysteresis and poor stabilized device efficiency.10

Addition of costly fullerene interlayers or fullereneterminated self-assembled monolayers (SAMs) between electron-selective MOx contacts and MAPbI3 PALs leads to significant improvements in PV device stability and performance.2, 10-12 Functionalization of MOx surfaces with cost-effective, bifunctional molecular modifiers, such as (3aminopropyl)triethoxysilane (APTES) and para-substituted aromatic carboxylic acids, has also been reported to improve MAPbI3 PV device performance.13-15 Substantial performance and stability enhancements have additionally been reported for hybrid perovskite PV devices incorporating chlorine-capped TiO2 nanocrystal interlayers.3 These improvements in performance and stability have been attributed to interface dipoles, passivation of oxide and perovskite defects and improved surface free energy (i.e., wettability). However, the fundamental chemical processes at the oxide interface and role of interlayers on the interfacial composition and energetics of PALs remain elusive due to the absence of concomitant characterization studies on bare and passivated oxide surfaces. In this study, we show that Brønsted and Lewis acid-base interactions between TiO2 surface defects and vacuum processed MAPbI3 precursors critically impact the growth, chemical composition and electronic structure of SILs. In situ photoelectron spectroscopy (PES) and mass spectroscopy (MS) analysis during film growth suggests that high coverages of methyl-terminated species (e.g., methylamine) adsorbed at TiO2 defect sites lead to (i) high surface free energy, (ii) islandlike growth of large and PbI2-rich nuclei, (iii) inhibited MAPbI3

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Figure 1. MAPbI3 surface species and passivation of TiO2 surface defects. (a) Bulk-truncated MAPbI3(100) structure with possible MAI and PbI2 terminations. (b) A device-relevant TiO2 surface is schematically represented using a bulk-truncated anatase TiO2(001) surface, showing defects associated with undercoordinated Ti and O atoms and ambient adsorbates. (c) Hydrolyzed APTES molecules passivate TiO2 surface defects via condensed silanol bonds with terminal OH groups (OHt,1c) and Lewis and Brønsted acid-base interactions between the terminal amine and TiO2 surface defects that lead to primarily horizontal rather than upright orientations. See text for additional details.

nucleation and (iv) energetic barriers for charge transport and extraction. Gas-phase modification of the oxide surface with multifunctional aminosilanes, which simultaneously passivate both Lewis and Brønsted defect sites on the TiO2 surface, (i) mitigates high coverages of methyl-terminated species, (ii) promotes conformal growth of significantly smaller PbI2-rich nuclei, (iii) enhances MAPbI3 nucleation and (iv) improves charge transport and extraction energetics. Elucidating and controlling interactions and reactions between PAL precursors and MOx defects, which are governed by unique Lewis16 and Brønsted17 acid-base chemistry associated with the identity of the metal cation (e.g., M = Ti, Sn, Ni, etc.) and precursors, is essential for further optimizing the efficiency and long-term stability of low-cost, scalable and efficient PV devices.6, 18

RESULTS AND DISCUSSION Defective and Aminosilane-Passivated TiO2 Surfaces. Proposed MAPbI3 and TiO2 surface species prior to contact are summarized using bulk-truncated structural schematics in Figure 1,19 which are based on cubic MAPbI320 and anatase TiO2.21 The TiO2 films investigated here are smooth (ca. 2-3 nm rms roughness) and amorphous,22-23 which implies a higher density of defects relative to polycrystalline films, and the bulktruncated anatase structure is simply used as a graphic illustration to facilitate discussion of typical near-surface defects and proposed APTES binding modes. The bulktruncated MAPbI3 structure provides MAI-rich and PbI2-rich terminations (Figure 1a). Theoretical models have suggested that the MAI-terminated surface is energetically more stable on anatase single crystal surfaces.24-25 However, we show here that interactions between MAPbI3 precursors and device-relevant TiO2 contacts result in a MA-deficient and PbI2-rich SIL prior to nucleation of MAPbI3. Based on established surface science studies,26 dangling bonds associated with undercoordinated and reactive Ti and O

atoms on the TiO2 surface lead to a number of potential adsorbates (Figure 1b). Oxidized Ti5c4+ sites and reduced Ti4c3+ species near oxygen vacancy (VO) and Ti interstitial (TiI) defects are strong Lewis acids, which form robust coordinate covalent bonds with strong Lewis bases (e.g., primary amines).16 Electron transfer from reduced defects to adsorbed molecular oxygen (O2(ads)) leads to superoxide (O2–•(ads)) species.27 Two unique pKa values are associated with terminal, 1-coordinate (OHt,1c; pKa ≈ 7.8) and bridging, 2-coordinate (OHbr,2c; pKa ≈ 5.0) surface hydroxyls.17, 28 Adventitious N2,29 carboxylate (e.g., formate and acetate)30 and aliphatic carbon species31 (not shown for clarity) may also occupy defect sites. Binding modes for hydrolyzed APTES are proposed in Figure 1c. Reaction of APTES with the oxygen plasma treated TiO2 surface (bare “TiO2” sample) under anhydrous, vaporphase conditions occurs via alcohol condensation with the most reactive terminal hydroxyl groups.32 APTES molecules can also physisorb to the TiO2 surface via hydrogen, ionic and coordinate covalent bonds with the terminal amine and unreacted alkoxy groups. Subsequent treatment in dilute HCl (“APTES-HCl” sample) or KOH (“APTES-KOH” sample) solutions hydrolyzes free ethoxy groups and changes the degree of amine protonation.33-34 While the aminopropyl end group of surface-bound APTES molecules is often assumed to extend away from the oxide surface (cf. references 13 and 14), our findings and those of others indicate that the majority of amine groups undergo strong intramolecular and intermolecular interactions that preclude upright orientations (Figure 1c). Murray and co-workers have shown that acid and base treatment, along with amidization reactions, of APTES-functionalized MOx surfaces only impacted a relatively small fraction (< 30%) of the available amine population, of which the majority were proposed to be inactive due to hydrogen and ionic bonds with unbound silanol groups (forming cyclic structures) and hydroxyls on the oxide

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Figure 2. Amine bonding interactions and general orientation of APTES. (a) AR-XPS N 1s CL spectra for TiO2, APTES-HCl and APTES-KOH samples at a more bulk-sensitive (0°) and surfacesensitive/grazing (60°) take-off angle (TOA, see inset in left panel). The low BE peak corresponds to electron-rich amines, and the high BE peak is associated with electron-deficient amines that interact strongly with surface defects (see inset in center panel). (b) Changes in N/Ti, Si/Ti and N/Si atomic ratio at a TOA of 60° w.r.t. 0° indicate that the APTES N atom is closer to the TiO2 than the Si atom and confirm that the amine plays a key role in passivating surface defects (see Figure 1c).

surface.34 Meroni et al. have recently reported plane-wave density functional theory (DFT) simulations, along with a suite of surface characterization techniques, that indicate more horizontal orientations for aminopropyl end groups due to hydrogen bonding interactions between the amine and TiO2 surface hydroxyls.35 In addition to these Brønsted acid-base and hydrogen bonding interactions, Farfan-Arribas and Madix have shown evidence that alkylamines (e.g., ethylamine) exclusively interact with undercoordinated Ti sites on stoichiometric and defective TiO2 surfaces via coordinate covalent bonds (ca. 2324 kcal/mol),36 which are significantly stronger than hydrogen bonds (< 10 kcal/mol). These previous theoretical and experimental investigations indicate that strong interactions between the APTES terminal amine and TiO2 surface defects result in a more horizontal orientation of the aminopropyl end group. The relative APTES coverage for the APTES-HCl and APTES-KOH sample is estimated to be ca. 40% and 50% of a close-packed (upright) monolayer, respectively (see Supporting Information, SI, Table S1 and Figure S1). Reduced APTES coverage (ca. 20%) after HCl treatment is explained by protonation and desorption of physisorbed molecules (see Figure 1c). These surface coverages correspond to a Si/Ti atomic ratio of 0.09 ± 0.01 and 0.13 ± 0.01 for APTES-HCl and APTES-KOH samples, respectively (Figure S2). HCl and KOH treatment also introduces low coverages of adsorbed Cl– and K+ ions (Cl/Si and K/Si atomic ratios ~ 0.2), but the Cl/Ti atomic

ratio after HCl treatment is ca. 7 times lower than the value reported for chlorine-capped TiO2 nanocrystal films.3 Angle-resolved XPS (AR-XPS) core level (CL) spectra, acquired at take-off angles (TOAs) of 0° (more bulk-sensitive) and 60° (more surface sensitive/grazing), are evaluated to determine the amine binding modes and general APTES orientation (Figure 2). Complete AR-XPS analysis of bare and APTES-modified TiO2 surfaces is provided in the SI (Figures S1-S2). An increase in the N 1s signal after APTES modification suggests at least two fitting components that are typically assigned to unprotonated (low BE) and protonated (high BE) amine species (Figure 2a).33 We suggest, however, that the N 1s envelope includes a variety of overlapping amine/ammonium species that increase in BE for more electron deficient nitrogen atoms. The low BE N 1s peak (NLBE) is assigned here to free base and H-bond donor, and the high BE peak (NHBE) is attributed to a range of electron-deficient species (e.g., H-bond acceptor, Lewis base and protonated), which dominate the N 1s spectra regardless of HCl/KOH treatment. A decrease in the relative population of electron-deficient NHBE species at grazing emission indicates their closer relative proximity to the oxide surface and is consistent with the terminal amine (pKa ≈ 10.6) abstracting protons from surface hydroxyls (pKa < 8) or bonding with undercoordinated Ti atoms.36 The average orientation of bound APTES is estimated from changes in the relative magnitude of N/Ti, Si/Ti and N/Si atomic ratios at a TOA of 60 ° with respect to (w.r.t.) 0° (Figure 2b). While the Si/Ti and N/Ti ratio both significantly increase at grazing emission, the considerably larger increase in Si/Ti ratio signifies that the amine/ammonium group is located below the Si atom, which is further supported by a ca. 20% decrease

Figure 3. Growth of co-evaporated PbI2/MAI films. (a,b) Attenuation of the Ti 2p3/2 (a) and O 1slat (b) signal during film growth on TiO2, APTES-HCl and APTES-KOH surfaces. The legend shows the measured and expected inelastic mean free path (ln) for LBL film growth. Hindered substrate attenuation on the bare TiO2 surface indicates that initial film growth is island-like, and close agreement with the expected attenuation indicates more conformal film growth on APTES-passivated TiO2 surfaces.


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in the N/Si ratio at grazing emission. These findings confirm a more horizontal aminopropyl orientation and indicate the terminal amine/ammonium group plays a key role in passivating strongly interacting and reactive TiO2 surface defects. Film Growth and Composition. MAI/PbI2 precursors are incrementally co-evaporated to yield films with nominal thickness ranging between 2 and 200 nm onto TiO2, APTESHCl and APTES-KOH samples.7 Formation of SILs and nucleation of stoichiometric MAPbI3 crystallites during cosublimation can occur through intermediate states that result from gas-phase MAI reactions, which include dissociation into the parent compounds (equation (1a)) or chemical transformation into ammonia and methyl iodide (equation (1b)):37 CH3NH3+I–(s) ⇌ CH3NH2(g)+HI(g) (1a) + – CH3NH3 I (s) ⇌ NH3(g)+CH3I(g) (1b) MS analysis in Figure S3 shows that equation (1a) is favored and suggests that adsorption of methylamine and HI at Ti or Pb Lewis acid sites (equation (2a)) and surface coupling (equation (2b)) takes place prior to reaction with PbI2-rich nuclei and MAPbI3 nucleation (equation (2c)): (2a) CH3NH2(g)+HI(g) ⇌ CH3NH2(ads)+HI(ads) CH3NH2(ads)+HI(ads) ⇌ CH3NH3+I–(ads) (2b) CH3NH3+I–(ads)+PbI2(ads) ⇌ MAPbI3(ads) (2c) Adsorbed HI may also dissociate with iodide anions and protons bonding to Lewis acid and Brønsted base sites, respectively. Therefore, desorption and side reactions of precursors at the oxide/PAL interface compete with film growth and MAPbI3 nucleation. Striking differences in the attenuation of Ti 2p and O 1s XPS CL signals for nominal film thicknesses less than 20 nm show that APTES passivation of the TiO2 surface significantly impacts the film growth mechanism (Figure 3 and Figure S4).

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The measured inelastic mean free path (ln,meas) of the Ti 2p3/2 (Figure 3a) and O 1slat (Figure 3b) signal is compared with the expected ln for layer-by-layer (LBL) film growth (ln,LBL).38 Agreement between the measured and expected ln values for APTES-modified TiO2 contacts indicates more conformal growth of substoichiometric nuclei and MAPbI3 crystallites. A ca. 3-fold increase in ln,meas for the film deposited on the bare TiO2 surface suggests island-like film growth with bare or sparsely covered regions of TiO2 present during the initial phases of deposition (see inset in Figure 3a). The TiO2 substrate signals are extinguished at a nominal film thickness of 20 nm, implying that the islands eventually coalesce into a more conformal film. We suggest that saturated coverages of methylterminated species, such as methylamine and adventitious carboxylate species,30, 36 result in a high free energy surface that leads to dewetting of polar PbI2-rich nuclei and island-like growth during the initial phases of precursor co-evaporation. APTES functional groups passivate TiO2 surface defects and mitigate high coverages of methyl-terminated molecules, and we propose that unbound silanol groups improve surface free energy and facilitate growth of more conformal seed nuclei and polycrystalline PAL films. PAL stoichiometry, which is reported as the atomic ratio w.r.t. Pb and extracted from XPS CL spectra in Figure S5, demonstrates that TiO2 surface chemistry drastically influences SIL composition and MAPbI3 nucleation (Figure 4a-4c). Significantly, the “nucleation thickness,” which is the nominal thickness where the atomic ratios reach the expected MAPbI3 stoichiometry, is 2-3´ thinner for the APTES-passivated TiO2 contacts. Consistent with previous reports,7-8 the SIL is deficient in N and, thus, MA near the oxide interface. A recent solution-based study has revealed that coordination of primary amines to PbI2 leads to Pb-alkylamide byproducts;39 however, absence of amide species in our N 1s spectra (Figure S5) suggests that amine desorption likely occurs prior to reaction in vacuum.36 Compared to a sharp increase in N that accompanies

Figure 4. PAL stoichiometry and energetic shifts from XPS. (a-c) Atomic ratios of I, N and PAL C (CPAL) w.r.t. Pb on TiO2 (a), APTESHCl (b) and APTES-KOH (c) samples. Prior to reaching the nucleation thickness (vertical dashed line), which is defined when atomic ratios reach the MAPbI3 stoichiometry, the substoichiometric interface layer (SIL) is MA-deficient and PbI2-rich. (d-f) XPS BE shifts for substratespecific (Ti 2p3/2 and O 1slat w.r.t. 0 nm) and precursor-specific (Pb 4f7/2, I 3d5/2 and N 1s w.r.t. 200 nm) CLs on TiO2 (d), APTES-HCl (e) and APTES-KOH (f) surfaces indicate charge transport barriers for PALs on TiO2 and APTES-HCl contacts. The trend lines in all plots are a guide to the eye.


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Figure 5. Film growth and MAPbI3 nucleation. (a) MAI primarily dissociates into methylamine and HI. (b) Methylamine and HI adsorb onto PbI2 and TiO2 surfaces, where high coverages of methylamine lead to dewetting of PbI2-rich nuclei on bare TiO2 and island-like growth prior to MAPbI3 nucleation. APTES passivates methylamine adsorption sites and provides a polar surface for growth of more conformal PbI2-rich nuclei. (c) Inhibited MAPbI3 nucleation on bare TiO2 is consistent with considerably longer diffusion lengths required for MAI coupling and subsequent reaction with large PbI2-rich nuclei with a smaller relative surfacearea-to-volume ratio.

nucleation on the bare TiO2 contact (Figure 4a), the relative N content steadily increases up to nucleation on the APTESmodified TiO2 contacts (Figure 4b and 4c), indicating a MA concentration gradient within the SIL. The relative concentration of N tracks a slight excess of PAL-related carbon (CPAL) and is interpreted as a surface layer of MA with the ammonium group pointed toward the film surface. Prior to measuring an observable N 1s signal (< 10 nm), the I content within the film on the bare TiO2 contact varies between ca. 2 and 2.5 in Fig 4a. The I content within the SIL on the APTES-HCl surface in Figure 4b is only slightly below 3 prior to nucleation when compared to a more PbI2-like stoichiometry on the APTES-KOH surface in Fig 4c. These observations imply that a higher relative surface concentration of unpassivated Lewis acid and protonated Brønsted acid species support higher interfacial concentrations of iodide species, which can migrate during PV operation and have been linked to J-V hysteresis and poor stabilized efficiency.6, 10 Proposed interactions that control the initial film growth mode, SIL composition and nucleation of MAPbI3 during coevaporation of MAI and PbI2 precursors near the TiO2 interface are summarized in Figure 5. In the gas-phase, MAI primarily dissociates into methylamine and hydroiodic acid (Figure 5a). Since MAI evaporation has been shown to be omnidirectional and alkylamines (i.e., methylamine) strongly bind to undercoordinated Ti sites,36, 40 the unpassivated TiO2 surface is quickly saturated with methyl-terminated molecules prior to PbI2 deposition and becomes non-polar, which results in islandlike growth of large, PbI2-rich nuclei (Figure 5b). Hydrolyzed

APTES molecules simultaneously passivate strongly interacting TiO2 defect sites and provide hydrophilic surface sites (e.g., unbound silanol groups) for more conformal growth of relatively small, PbI2-rich nuclei (Figure 5b). MAPbI3 nucleation, which we suggest is preceded by surface coupling of methylamine and HI to form MAI, competes with desorption and/or dissociation of these precursors (Figure 5c). We propose that inhibited MAPbI3 nucleation on the unpassivated TiO2 surface can be explained by longer diffusion lengths (LD) for MAI coupling and subsequent reaction with large, PbI2-rich nuclei, and enhanced MAPbI3 nucleation on APTESfunctionalized TiO2 surfaces is consistent with shorter LD for MAI coupling and MAPbI3 nucleation with smaller, PbI2-rich nuclei (Figure 5c). This qualitative growth and nucleation model assumes that TiO2 surface chemistry influences the average nuclei size and, thus, surface-area-to-volume ratio, which is larger for smaller nuclei and decreases the relative probability for desorption of volatile precursors prior to reaction and nucleation of MAPbI3. X-ray diffraction (XRD) analysis provides evidence that TiO2 surface defects also considerably impact the composition of solution-processed MAPbI3 and MABr-doped FAPbI3 PAL films (ca. 400 nm) with systematically controlled organohalide salt concentrations (Figure S6). In general, an increase in the organohalide salt concentration leads to a decrease in the relative concentration of unreacted PbI2, which is driven to interfaces (i.e., grain boundaries and surfaces) during PAL crystallization.23 However, significant levels of residual PbI2 persist within PAL films at the highest salt concentration on unpassivated TiO2 contacts and suggest that Lewis and Brønsted acid-base interactions and reactions between the organohalide salt and TiO2 surface defects produce impurities that inhibit PAL nucleation. These decomposition products, which are also driven to interfaces during PAL crystallization, are likely responsible for the significantly reduced coherence length of residual PbI2 crystallites on poorly passivated TiO2 and incompletely passivated APTES-HCl surfaces (Figure S6). Energetics of TiO2/MAPbI3 Heterojunctions. The impact of surface passivation on the electronic structure of TiO2/PAL heterojunctions is evaluated by analyzing XPS and ultraviolet photoelectron spectroscopy (UPS). Changes in the band edge energies of the buried TiO2 contact during film growth are initially evaluated by analyzing the O 1slat and Ti 2p3/2 CL BE shifts w.r.t. the pristine surface (0 nm) in Figure 4d-4f. The absolute BEs can be found in Figure S7. The O 1slat and Ti 2p3/2 BE increases and equilibrates at the same value for all three TiO2 samples, implying that the conduction band minimum energy (ECBM) is pinned at the same energy. This energetic equilibration is consistent with enhanced rates of superoxide attenuation during film growth (Figure S4), suggesting superoxide desorption and concomitant reduction of TiO2 relaxes initial interfacial depletion conditions (equation (3)). Ti4c4+–O2•–(ads) → Ti4c3++O2(g) (3) Equilibration of TiO2 energetics is further interrogated by correlating the valence band maximum energy (EVBM) from UPS (Figure S8) with initial XPS BE positions and subsequent shifts during film growth (Figure S9). This analysis strongly suggests that ECBM is pinned just above EF at the interface after equilibration, which we assume leads to flat band conditions. Relaxation of the interface depletion layer is accompanied by reappearance of a localized VO/Ti4c3+ gap state (GSVo) at ca. 1.2 eV below EF in the UPS spectra near the interface (Figure S10). These gap states can improve the photoconductivity of TiO2 and


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Figure 6. Energy level diagrams. Co-evaporated MAI/PbI2 films interfaced with TiO2 (a) APTES-HCl (b) and APTES-KOH (c) contacts. The MA-deficient and PbI2-rich SIL is shown as a yellow/hatched region. See text for additional details.

lead to enhanced charge extraction, performance and stability of hybrid perovskite PV devices.41 BE shifts for PAL CLs are determined w.r.t. the “bulk” MAPbI3 film (200 nm) and are shown in Figure 4d-4f. The Pb 4f7/2 and I 3d5/2 CL peaks shift and equilibrate at bulk-like BEs (Figure S6).7 For the TiO2 (Figure 4d) and APTES-HCl (Figure 4e) samples, which are associated with enhanced concentrations of interfacial iodide species (Figure 4a and 4b), the Pb 4f7/2 and I 3d5/2 BEs shift above the bulk BE and imply the presence of a charge transport barrier. Conversely, the wellpassivated APTES-KOH sample, which shows the lowest relative excess of interfacial iodide species (Figure 4c), facilitates a smooth upward shift in the Pb 4f7/2 and I 3d5/2 BE that suggests ideal band bending for charge transport within the SIL and MAPbI3 film (Figure 4f). We suggest that orthogonal N 1s BE shifts indicate a reduced positive charge on the MA nitrogen atom due to hydrogen bonds with excess interfacial iodide species. UPS analysis during film growth provides more detailed insight into the frontier orbital energetics of the PAL (Figure S11), which are summarized, along with the equilibrated TiO2 energetics (Figures S7-S10), using energy level diagrams (Figure 6). The precision of the reported energies from UPS analysis of the PALs in Figure 6 is ≤ 0.04 eV. These diagrams provide insight into charge transport within the MAPbI3 film and electron collection at the TiO2 contact, both of which are critical factors that control PV performance and stability. The SIL is highlighted by a yellow/hatched area in Figure 6. An abrupt decrease in EVBM coincides with MAPbI3 nucleation on the bare TiO2 sample in Figure 6a. A constant EVBM between the SIL and the film bulk indicates the absence of band bending on the bare TiO2 contact (DEVBM,TiO2 = 0 eV), suggesting the thick SIL leads to poor electronic coupling between the TiO2 contact and MAPbI3 film.25 A similar degree of band bending (DEVBM,APTES-HCl = 0.20 eV and DEVBM,APTES-KOH = 0.24 eV) is observed for the films on the APTES-treated TiO2 contacts in Figure 6b and 6c; these shifts asymptote at a nominal thickness of 50 nm, which approximates the width of the accumulation layer. Agreement between the thickness-dependent N

composition and EVBM confirms that band bending is enabled by the MA concentration gradient (Figure S13).9 Bulk EVBM and work function (F) values are in close agreement with previous measurements of MAPbI3 films on n-type MOx contacts.42-43 We note the absence of significant charge transport barriers from the UPS-derived EVBM values used to construct the energy level diagrams (Figure 6) that are inferred from CL BE shifts for PAL films on TiO2 and APTES-HCl samples in Figure 4. This discrepancy is likely related to the difference in probe depth for XPS (dp,XPS ~ 10 nm) and UPS (dp,UPS ~ 2 nm) measurements.38 Nevertheless, XPS and UPS analysis clearly indicate that the well-passivated APTES-KOH sample leads to the most beneficial energetics for charge transport within the PAL. ECBM is estimated for the MAPbI3 layer by addition of the optical gap (Eg,opt ≈ 1.6 eV) to EVBM in Figure 6. For thicknesses that show significant N deficiencies within the SIL, ECBM is estimated by addition of the PbI2 optical gap (Eg,opt ≈ 2.2 eV).43 This PbI2-rich interface layer introduces a ca. 0.6 eV energy barrier for electron transfer from MAPbI3 to the TiO2 contact. We suggest that thick and MA-deficient SILs, which result from uncontrolled interface chemistry, are responsible for the previously reported ECBM mismatch between amorphous TiO2 contacts and MAPbI3.18 The difference between changes in the bulk work function (DF) and band bending across the interface yields the total interface dipole (eDtot, equation (4)) reported in Figure 6:44 eDtot = DF – DEVBM (4) The unmodified TiO2/MAPbI3 heterojunction yields the largest interface dipole (eDTiO2(tot) = 0.40 eV), which compensates for the absence of band bending in the active layer. Due to enhanced band bending and a smaller bulk work function, the smallest interface dipole is found on the APTES-KOH contact (eDAPTES-KOH(tot) = 0.14 eV), and the interface dipole is slightly larger for the APTES-KOH/MAPbI3 heterojunction (eDAPTESHCl(tot) = 0.20 eV). Therefore, smaller interface dipoles indicate decreased contact reactivity and improved interfacial energetics.45


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The energy level diagrams in Figure 6 demonstrate that APTES passivation of TiO2 surface defects leads to thinner and compositionally-graded SILs and improved interfacial energetics that control the efficiency of charge transport and collection. These studies show that optimization of PAL/MOx heterojunction composition and electronic structure is enabled by understanding and controlling the surface composition and chemistry of low-cost MOx electrical contacts using costefficient and scalable APTES SAMs.

CONCLUSIONS This work elucidates the critical impact of TiO2 surface defects on the growth, composition and energetics of coevaporated MAPbI3 films. We have demonstrated an analytical approach to fully characterize TiO2 surface defects and passivation strategies that can be extrapolated to a much wider range of oxide interlayers in these emerging PV technologies. Lewis and Brønsted acid-base interactions between PAL precursors and electron- and hole-collecting MOx electrical contacts will uniquely depend on the processing conditions, precursor chemistry and identity of the metal cation (e.g., Ti, Sn, Ni, etc.), which are expected to influence the resultant composition and energetics of the PAL. Low-cost and scalable passivation strategies that target specific chemical reactivities between these oxides and PALs are needed, and we have shown that multifunctional aminosilanes simultaneously passivate multiple TiO2 surface defects and mitigate chemical interactions that are expected to negatively influence the longterm stability and performance of PALs in low-cost and efficient PV device technologies.

EXPERIMENTAL SECTION Preparation of TiO2 and APTES-modified TiO2 Thin Films. Compact and amorphous TiO2 films (ca. 30 nm thick, rms roughness ca. 2-3 nm) are deposited onto oxygen plasma treated indium tin oxide (ITO) substrates in a home-built chemical vapor deposition (CVD) system that has been described in a previous publication.22-23 Bare “TiO2” samples are treated with oxygen plasma (17 W, ca. 800 mTorr, 10 min) to remove adsorbed surface contaminants. APTES is adsorbed from the vapor-phase to TiO2 samples in a N2 glovebox (< 0.1 ppm H2O and < 1 ppm O2) using a previously-reported procedure with some modifications.46 The samples are transferred into ambient, rinsed with toluene and ethanol, dried with N2 and transferred back into the N2 glovebox where they are soft-baked on a hotplate at 120 °C for 10 min. Prior to loading the samples into the ultra-high vacuum (UHV) system, the aminosilane-treated TiO2 substrates are dipped into freshlyprepared HCl (50 mM, pH » 1.3) or KOH (2 mM, pH » 11.3) solutions for 15 s, which has yielded consistent and reproducible results, and dried with a stream of N2 to yield “APTES-HCl” or “APTES-KOH” samples, respectively. PES Measurements. The bare TiO2 and aminosilane-modified TiO2 samples are mounted in air and loaded into a UHV system for PES measurements. XPS (monochromatic Al Ka excitation at 1486.6 eV, 300 W, pass energy of 20 eV) measurements of the pristine substrates are acquired using a Kratos Axis Ultra PES system with a base pressure of ca. 2 x 10-9 Torr. For XPS throughout the vacuum co-evaporation experiments, CL spectra are taken at a takeoff angle of 0° using a nonmonochromatic Mg Ka XPS source (1253.6 eV, 12 kV, 20 mA) and a Phoibos 100 hemispherical analyzer (pass energy of 10 eV). UPS measurements are taken with a monochromatic VUV 5000 microwave UV source (VG Scienta) using the He Ia emission line (21.22 eV) with a -8 V sample bias and an analyzer pass energy of 2 eV. We did not observe any change in the BE or shape of the XPS CL or UPS VB/SECO spectra for any of the samples during analysis with UV or X-ray irradiation. Co-evaporation of MAI and PbI2. The thermal co-evaporation of MAPbI3 precursor films onto the TiO2 contacts and subsequent PES

characterization have been conducted over a two-week period and the procedures are very similar to a previous report.7 Briefly, MAI and PbI2 are evaporated from separate quartz Knudsen cells in the growth chamber (base pressure in the mid 10-8 mbar range) at rates of 0.60 Å/s (ca. 120 °C) and 0.40 Å/s (ca. 300 °C), respectively, which are measured with individually-calibrated QCMs, at a pressure of ca. 4´10-4 mbar. After each deposition, the films are transferred without breaking vacuum to the preparation chamber (base pressure in the low 10-9 mbar range) and annealed at 70 °C for 1 hour. After cooling, the samples are transferred to the analysis chamber (high 10-10 mbar range) for PES measurements (additional experimental and measurement details are provided in the Supporting Information, SI).

ASSOCIATED CONTENT SUPPORTING INFORMATION The Supporting Information is available free of charge on the ACS Publications website. Additional experimental details, XPS CL and UPS VB/SECO spectra and analysis for co-evaporated films and TiO2 contacts, MS analysis during co-evaporation of MAI/PbI2 and XRD patterns and analysis for solution-processed PALs.

AUTHOR INFORMATION Corresponding Author *Email: [email protected]

Notes

The authors declare no competing financial interests.

ACKNOWLEDGEMENTS This research was partially supported by the National Science Foundation (DMR: 1506504 and DMR: 1337371) and a grant from the Arizona Board of Regents, Research Innovation Fund. Current and future support for this project is provided by the Office of Naval Research (ONR: N00014-18-1-2711). RCS acknowledges salary support from the Office for Research, Discovery and Innovation (RDI) at the University of Arizona and travel support from the University of Cologne.

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