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Dedoping of Lead Halide Perovskites Incorporating Monovalent Cations Mojtaba Abdi-Jalebi, Meysam Pazoki, Bertrand Philippe, M. Ibrahim Dar, Mejd Alsari, Aditya Sadhanala, Giorgio Divitini, Roghayeh Imani, Samuele Lilliu, Jolla Kullgren, Håkan Rensmo, Michael Grätzel, and Richard H. Friend ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b03586 • Publication Date (Web): 28 Jun 2018 Downloaded from http://pubs.acs.org on June 29, 2018

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NaI

CuI

AgI

Valence edge

Intensity

Pristine

Valence Band (VB)

CB

VB

CB

CB

Binding Energy (eV)

Fermi Level

VB

VB

Pristine NaI CuI AgI

(110)

Intensity

Conduction Band (CB)

Energy

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Dedoping of Lead Halide Perovskites Incorporating Monovalent Cations Mojtaba Abdi-Jalebi1*, Meysam Pazoki2,3, Bertrand Philippe4, M. Ibrahim Dar5, Mejd Alsari1, Aditya Sadhanala1, Giorgio Divitini6, Roghayeh Imani3, Samuele Lilliu7,8, Jolla Kullgren3, Håkan Rensmo4, Michael Grätzel5, Richard H. Friend1 1

Cavendish Laboratory, Department of Physics, University of Cambridge, JJ Thomson Avenue, Cambridge CB3 0HE, UK.

2

Department of Engineering Sciences, Solid State Physics, Uppsala University, Box 534, SE 751 21 Uppsala, Sweden

3

Department of Chemistry, Ångström Laboratory, Uppsala University, Box 538, SE 75121 Uppsala, Sweden.

4

Molecular and Condensed Matter Physics, Department of Physics and Astronomy, Uppsala University, Box 516, SE 75120

Uppsala, Sweden. 5

Laboratory of Photonics and Interfaces, Institute of Chemical Sciences and Engineering, École Polytechnique Fédérale de

Lausanne, Lausanne, CH-1015, Switzerland. 6

Department of Materials Science & Metallurgy, University of Cambridge, 27 Charles Babbage Road, Cambridge CB3 0FS,

UK 7

Department of Physics and Astronomy, University of Sheffield, Sheffield, S3 7RH, UK.

8

The UAE Centre for Crystallography, UAE.

*[email protected]

ABSTRACT: We report significant improvements in the optoelectronic properties of lead halide perovskites with the addition of monovalent ions with ionic radii close to Pb2+. We investigate the chemical distribution and electronic structure of solution processed CH3NH3PbI3 perovskite structures containing Na+, Cu+ and Ag+, which are lower valence metal ions than Pb2+ but have similar ionic radii. Synchrotron x-ray diffraction reveals a pronounced shift in the main perovskite peaks for the monovalent cation-based films, suggesting incorporation of these cations into the perovskite lattice, as well as a preferential crystal growth in Ag+ containing perovskite structures. Furthermore, the synchrotron X-ray photoelectron measurements show a significant change in the valence band position for Cu- and Ag-doped films although the perovskite bandgap remains the same, indicating a shift in the Fermi level position towards the middle of the bandgap. Such a shift infers that incorporation of these monovalent cations de-dope the ntype perovskite films when formed without added cations. This de-doping effect leads to cleaner bandgaps as reflected by the lower energetic disorder in the monovalent cation doped perovskite thin films as compared to pristine films. We also find that in contrast to Ag and Cu, Na locates mainly at the grain boundaries and surfaces. Our theoretical calculations confirm the observed shifts in x-ray diffraction peaks and Fermi level, as well as absence of intra-bandgap states upon energetically favorable doping of perovskite lattice by the monovalent cations. We also model a significant change in the local structure, chemical bonding of metal-halide and the electronic structure in the doped perovskites. In summary, our work highlights the local chemistry and influence of monovalent cation dopants on crystallization and the electronic structure in the doped perovskite thin films. Keywords: monovalent cations; de-doped perovskite thin films; enhanced optoelectronic quality; substitutional doping; interstitial doping Page 1 of 32 ACS Paragon Plus Environment

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Hybrid organic-inorganic lead halide perovskites display significant intrinsic properties including high absorption coefficient,1 sharp and tunable bandedge,2,3 long charge carrier diffusion length,4,5 low trap densities, high luminescence quantum yields,6,7 and the photon recycling capability.8 These materials have shown a relatively fast evolution in the performance of both light emitting diodes and solar cells (e.g. power conversion efficiency of 22.7%) since their first appearance in the devices.9,10 Perovskite materials are characterised by the general formula of ABX3 including a single cation or a mixture of monovalent cations A (e.g. Cs, Rb, CH3NH3 (MA), CH(NH2)2 (FA)), a divalent metal B (e.g. Pb, Sn), and a halide X (e.g. Cl, Br, I). Compositional engineering of perovskite is key for customising the electronic properties of perovskite materials to obtain different optical bandgaps as well as higher device performance and stability.11 Recent calculations have shown that the electronic band structure is mainly affected by the exterior orbitals of the divalent cation and halide, and hence the monovalent cation does not affect the bonding directly12–14 although it can create indirect steric effects15 and interact with defects.16 As a result, a complete or partial substitution of the divalent metal can be used for engineering the light absorption coefficient, the band gap, and the charge carrier diffusion length.17,18 It has been shown that tuning the optoelectronic properties of perovskites is possible by inclusion or partially substituting Pb with another divalent cation (e.g. Sn, Sr, Cd and Ca),3,19 a monovalent cation (e.g. Na, Cu, Ag),20,21 or a trivalent cation (e.g. Bi, In, Au).22 This has led to further improvements in solar cell efficiency by adding K,23 Na,24–27 Cu28 or Ag29,30 into different perovskite systems. Furthermore, recent literatures show that substitution of two Pb cations with one monovalent cation (e.g. Ag) and one trivalent such as In or Bi cations can form double perovskites confirming the possibility of substitutional doping in metal halide perovskite structures.31–33 Herein, we explore the chemical distribution, chemical bonding, electronic and structural effects of Na, Cu and Ag monovalent cation iodide dopants in hybrid organic-inorganic perovskite films with enhanced optoelectronic quality upon doping. These monovalent cations have ionic radii of 116 pm Na, 91 pm Cu and 129 pm Ag, and can potentially replace Pb (133 pm) within the perovskite lattice34 or occupy the interstitial sites as observed for potassium ion.35 Notably, these monovalent cations cannot occupy the A site in the ABX3 perovskite structure since they are too small and mismatched

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for the A site in the lattice to sustain a photoactive ABX3 with an appropriate Goldschmidt tolerance factor.36 We investigate the effect of incorporating these dopants on the crystal growth of PbI2 and CH3NH3PbI3 perovskite with Grazing-Incidence Wide-Angle X-ray Scattering (GI-WAXS). We use soft X-ray Photoelectron Spectroscopy (XPS) and synchrotron Hard X-ray Photoelectron Spectroscopy (HAXPES) using synchrotron radiation to detect the cations and analyse the chemical composition and electronic structure of the perovskite from surface to bulk. We also performed Scanning Transmission Electron Microscopy-Energy Dispersive X-ray Spectroscopy (STEM-EDX) to probe the local chemistry of doped-perovskite thin films. We then perform a comprehensive Density Functional Theory (DFT) calculation to explore how the incorporation of these dopants in the perovskite lattice affects chemical bonding, atomic rearrangement in a unit cell, charge densities, and band structure. Our findings reveal a clear picture of the distribution of monovalent cation dopants and their effects on the structure and growth of perovskite films. RESULTS AND DISCUSSION We processed CH3NH3PbI3 perovskite films doped with Na, Cu and Ag monovalent cation dopants using solution-based sequential deposition technique.37 We also used single step deposition technique to confirm that the incorporation of the monovalent cation dopants is independent of the processing methods (see Supporting Information (SI)). Absorption spectra from Photo-thermal Deflection Spectroscopy (PDS) (Figure S1) indicate that the band gap remains unperturbed upon doping. In Cudoped perovskites we observe an absorption tail originating from the intrinsic absorption of copper,20 while in Ag- and Na-doped perovskites we measure the lowest sub-bandgap absorption and the lowest electronic disorder. In Figure 1, we summarize the impact of these dopants on the optoelectronic quality of the perovskite thin films. Time-resolved photoluminescence experiments show significant enhancement in the carrier lifetime for the doped samples reaching the longest lifetime for Na- and Ag-doped perovskite thin films due to elimination of fast non-radiative component seen in the pristine sample (Figure 1a). In Figure 1b, we show the photoluminescence quantum efficiency (PLQE) of the pristine and doped perovskite thin films. The PLQE shows a significant enhancement from the value of 3% for the Page 3 of 32 ACS Paragon Plus Environment

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pristine film to 12% and 15% for Cu- and Ag-doped perovskite films, respectively, and reach the highest value of 19% for Na-doped sample. The enhanced PLQE and PL lifetime collectively suggest that these dopants are able to eliminate non-radiative decay channels in the perovskite thin films.

Figure 1. Enhanced optoelectronic quality of perovskite thin film via doping. (a) Time-resolved photoluminescence decays of the doped-perovskite films, with excitation at 400 nm and a pulse fluence of 0.5 µJ cm−2. (b) PLQE of the doped perovskite films measured under illumination with a 532-nm laser. The inset shows the PL spectrum of the pristine and doped perovskite films. (c) Opencircuit voltage (Voc) and (d) short-circuit current (JSC), (e) fill factor (FF) and (f) power conversion efficiency (PCE) of the subsequent devices with standard deviation (SD) extracted from currentvoltage characteristics of pristine and doped-perovskite devices measured under full simulated solar illumination conditions (AM1.5, 100 mW.cm-2) and scanned at a rate of 15 mV.s-1 (see Figure S2 for J-V curves).

To validate how these luminescence results translate to device performance, we also fabricated perovskite solar cells using the device architecture fluorinated-tin oxide (FTO) / compact-TiO2 (~30 Page 4 of 32 ACS Paragon Plus Environment

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nm) / thin-mesoporous TiO2 (~200 nm) / perovskite (~500 nm) / Spiro-OMeTAD (~150 nm) / Au (80 nm). We summarise the changes in the photovoltaic parameters measured under full simulated sunlight for pristine and doped-perovskite based solar cells in Figure 1c-f (see Figure S2 and Table S1, for current-voltage (J-V) curves and device parameters, respectively). As expected from the relationship between luminescence and open-circuit voltage, the changes in open-circuit voltage (Voc) of the devices closely follow the PLQE and PL lifetime trends (Figure 1a-c). Furthermore, we see an increase in the short circuit current (Jsc) and fill factor (FF) of the doped-perovskite based devices (Figure 1d, e) inferring relatively high quality of doped-perovskite films. Finally, the enhanced optoelectronic quality of perovskite thin films via doping reflected in the power conversion efficiency (PCE) of the devices reaching 20.5%, 19.2% and 20.0% for Na-, Cu- and Ag-doped solar cells, respectively, compared to 16.7% for the pristine one. Next, we investigate the influence of these dopants on the growth of PbI2 with synchrotron Grazing Incidence Wide Angle X-ray Scattering38 (GI-WAXS) measurements (Figure S3) performed at the XMaS beamline (ESRF, France). With the Ag dopant (Figure S3d), the hexagonal PbI2 crystals are preferentially oriented with the c lattice vector perpendicular to the substrate (out-of-plane orientation), as is clear from azimuthal distribution of the (001) PbI2 ring at q ≈ 0.9 Å-1 in the GIWAXS pattern. When Na and Cu are added (Figure S3b-c), the PbI2 crystals display patterns similar to that of the pristine PbI2 (Figure S3a), i.e. a broader azimuthal distribution of the (001) PbI2 ring, indicating that a certain amount of crystals are randomly oriented with respect to the substrate. Plots of integrated radial line profiles (intensity vs azimuthal angle χ) are used to compare the relative orientation of the PbI2 crystals of the four samples (Figure S3e). Interestingly, the incorporation of Na, Cu, and Ag dopants results in an expansion of the PbI2 unit-cell along the c-axis, as evident from the shift of the (001) peak in the in-plane (Figure S3f) and out-of-plane (Figure S3g) azimuthally integrated line profiles.


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Figure 2. GI-WAXS diffraction patterns of (a) pristine, (b) Na-, (c) Cu- and (d) Ag-doped perovskite films. a, PbI2 indicates the (001) reflection of lead iodide and (002), (110) indicates the main perovskite diffraction ring, χ indicates the azimuthal angle. The molar ratio of the dopants in the precursor solution is 10 mol% with respect to Pb.

In Figure 2 a-d, we compare GI-WAXS diffraction patterns of pristine and doped-CH3NH3PbI3 (from conversion of PbI2 into perovskite using sequential deposition technique) measured with a laboratory beamline (XEUSS, XENOCS). We note a shift in the main perovskite peak (110) at q ≈ 1 Å-1 in the in-plane azimuthally integrated line profiles for the doped-perovskites (Figure S4), indicating that the aforementioned expansion of the PbI2 unit-cell along the c-axis seems to take place in the doped perovskite unit-cell as well, indicating the possibility of substitutional or interstitial doping with these monovalent cations. Though substitutional doping is expected to shrink the lattice, the lattice expansion does not preclude it as the chemical bonding and lattice rearrangement might be responsible for the lattice expansion (see the discussion in the calculations section). Furthermore, the combination of interstitial and substitutional can be a plausible scenario in which interstitial doping can cause the expansion of perovskite lattice due to the potential creation of stress/strain. In the Agdoped perovskite system, we observe a strong preferential orientation at χ ≈ 54° and a weaker preferential orientation at χ < 10° in the main perovskite ring (Figure 2d and Figure S4d). The other films (Na-doped, Cu-doped, and pristine perovskites) display a more uniform azimuthal distribution


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in the main perovskite ring, with some preferential orientation at intermediate angles (40° < χ < 60°) and at χ < 10° (Figure 2a-c and Figure S4). This suggests that the incorporation of Ag as a dopant in the precursor solution has a strong influence on the growth and final orientation of both PbI2 and perovskite crystallites. Notably, Ag-doping indicate a significant impact on the film growth as the formation of perovskites is slowed down (observed by the rate of colour change in the perovskite film) resulting in smoother film with more homogenously oriented grains as apparent from the topview SEM images (Figure S5). It is reported that slower film growth leads to higher degree of crystallographic orientation.39 Furthermore, the dopants can preferentially changes the energy of certain crystallographic planes in the crystal lattice associated with the size and overall potential of the cation. These collectively suggest that the origin of preferential origination that we observed in silverdoped PbI2 and perovskite thin films can be attributed to the amendment of the nucleation and growth rate of perovskite structures via Ag-doping. To get a deeper insight into the chemical composition, distribution, and electronic structure of the perovskite films at the surface and bulk we use soft X-ray Photoelectron Spectroscopy (XPS, 1486.6 eV photon energy) and synchrotron Hard X-ray Photoelectron Spectroscopy (HAXPES, 4000 eV photon energy) at the KMC-1 beamline (BESSY II, Germany), respectively. We show the overview spectra of the perovskite films recorded with two different energies in Figure S6 and the Pb 5d and I 4d spectra in Figure 3a-d. We observe a significant shift towards lower photon energies in the Pb5d and I4d spectra for the Ag- and Cu-doped perovskite films that overlap in binding energy. This shift is present in both the XPS and HAXPES spectra but it is larger in the latter. We also find a similar trend for the C 1s and N 1s core level (Figure S7) as well as for the valence band (VB) edge (Figure 3e-f). The observed shift in all the core level peaks as well as the valence band region is attributed to a modification of the Fermi level position within the bandgap of the doped-perovskites (e.g. Cu and Ag) as illustrated in Figure 3g indicating dedoping of perovskite thin films via incorporation of Cu and Ag. It is notable that we calibrated the spectra vs. the Fermi level and we detect a shift in all the core level peaks. To further confirm the observed dedoping effect in the Cu- and Ag-incorporated perovskite thin films, we performed conductivity measurements. In Figure S8, we show that the


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conductivity of the pristine sample (68 nS.m-1) dropped by an order of magnitude for Cu-doped (2 nS.m-1) and Ag-doped (6 nS.m-1) perovskites as a consequence of the dedoping effect. Notably, the position of Fermi energy within the bandgap of perovskite strongly depends on the processing methods, stoichiometry and the precursor compositions as well as the charge transport contacts. The highly n-doped pristine perovskite that we used in this work is consistent with previous reports using a similar approach for processing and device contacts.40,41


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Figure 3. (a)(b) Pb 5d, (c)(d) I 4d core level spectra and (e)(f) valence band edge of pristine and doped-perovskite thin films measured with a photon energy of 1486.6 eV (left panels) and 4000 eV (right panels). The I4d and Pb5d spectra were normalized versus the Pb5d5/2 core level. The valence spectra were normalized versus their maximum. See full valence spectra between -1 and 14 eV binding energy in Figure S8. (g) Schematic representation of the valence and conduction band edges position vs. the Fermi level (black dot line) induced by the incorporation of dopants.

In Figure 4, we measure the core level spectra of the dopants (i.e. Na 2s, Cu 2p and Ag 3d) in both photon energies (1486.6 and 4000 eV). We note that the Na 1s peak was detectable only in the XPS spectra but not in the HAXPES due to strong overlap with the iodine 3s core level (Figure S10). Despite introducing small dopant quantity, the overlap of the lines with more intense core levels (e.g. Cu 2p with I 3p, Na 1s with I 3s) and the low photoionization cross section of the core levels (e.g. Na 2s, Na 2p) particularly at 4000 eV, we detect all the monovalent cations introduced in the precursor solution by comparing the doped-perovskite spectra with the pristine one.


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Figure 4. (a)(b) Na2s, (c)(d) Cu2p and (e)(f) Ag3d core level spectra of the different perovskite films recorded with a photon energy of 1486.6 eV (top panels) and 4000 eV (bottom panels).

In Table 1, we summarize the ratio between the main elements in the perovskite films and the degree of dopants’ inclusion based on the intensity ratios between different atomic core levels acquired at 1486.6 eV and 4000 eV photon energies. We calculated the precursor composition ratios based on the concentration of the precursors. We also show the experimental I/Pb ratio in Table 1(a), which is comparable to the precursor composition ratio of 3 for the pristine and 3.1 for the doped-perovskite films. The content of iodide detected in the bulk (4000 eV) is very close to the precursor composition ratio, however, at the surface (1486.6 eV) an excess of iodide can be detected.42 We also notice that the content of iodide is slightly higher for the doped-films due to the addition of iodide from the initial salts (NaI, CuI and AgI). The excess of halide can fill halide vacancies and passivates traps.43

Table1. Intensity ratios between different core levels calculated from HAXPES measurements; (a) I/Pb ratio, (b) Na/Pb, Cu/Pb and Ag/Pb ratio for the different doped-perovskites as a function of the excitation energy (1486.6 and 4000 eV). The values for the precursor composition ratio correspond to the calculated ratios within the precursor solution. The error bar for these ratios is about ±0.01.

(a) I/Pb intensity ratio (precursor composition)

Pristine

Na

Cu

Ag

(3.0)

(3.1)

(3.1)

(3.1)

1486.6 eV

3.18

3.29

3.21

3.30

4000 eV

3.00

3.07

3.08

3.05

(b) Intensity ratio

Na/Pb

Cu/Pb

Ag/Pb

(precursor composition)

(0.1)

(0.1)

(0.1)

1486.6 eV

0.126

0.009

0.029

4000 eV

0.097

0.051

0.012


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In Table 1b, we also list the proportion of Na, Cu and Ag with respect to Pb based on the ratios between their core level peaks. These values suggest that Na has the highest inclusion into the perovskite film compared to the other dopants and has higher concentration at the top surface compared to deeper in the bulk. On the other hand, Cu penetrates more towards the bulk while traces of Ag can be detected more at the surface. The presence of Na does not affect the electronic structure of perovskites due to the complete overlap of its valence band with the pristine film. As for Cu and Ag, we find a lower inclusion of the monovalent cation compared to the precursor composition ratio of 0.1 (i.e. ratio of monovalent cation to Pb) but, surprisingly, this smaller quantity of Cu and Ag leads to dedoping and a stronger impact on the valence and conduction band energy positions (Figure 3g). To investigate the local chemical and morphological composition of the doped-CH3NH3PbI3 thin films, we performed Scanning Transmission Electron Microscopy-Energy Dispersive X-ray spectroscopy (STEM-EDX).44 In Figure 5, we show a stack of STEM High Angle Annular Dark Field (HAADF) cross sectional views of lamellae from the Na, Ag and Cu doped perovskite films. To preserve the perovskite layer during specimen preparation, we deposited capping layers of 2,2′,7,7′tetrakis(N,N-di-p-methoxyphenylamine)-9,9′-spirobifluorene (i.e. Spiro-OMeTAD, spin-casted) and platinum (ion-beam deposited), respectively (Figure 5a). From the STEM-EDX elemental analysis, we observe silver-rich inclusions well distributed within some of the perovskite grains (Figure 5b) while we find sodium-rich phases mostly at the grain boundaries of the perovskite as well as the interface with the substrate (Figure 5c). To study the Cu-doped thin film, the stray Cu signal, has been limited by the use of a molybdenum TEM grid and a low-background TEM holder, but the overall signal level makes isolating signal from the incorporated Cu in the perovskite layer very challenging (Figure 5d). In Figure 5e, f, we observe a relatively homogeneous distribution of lead and iodine within the perovskite layer across the entire specimen area. We also analysed the STEM-EDX dataset using multivariate analysis methods, specifically a Non-negative Matrix Factorisation (NMF) algorithm,45 which indicates the presence of a similar perovskite phase for all of the doped perovskite films (Figure S11).


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Figure 5. (a) Stack of HAADF STEM cross sectional image of perovskite thin film deposited on silicon, coated with Spiro-OMeTAD, and capped by Platinum layer. Elemental maps for (b) Ag, (c) Na, (d) Cu, (e) Pb and (f) I. For each element, we show Na, Ag and Cu- doped perovskite films from left to right. See the intensity spectra acquired from NMF decomposition dataset for each dopants in Figure S10. The scale bar is the same from (b) to (f).

Furthermore, we performed Scanning Electron Microscopy- Energy Dispersive X-ray spectroscopy (SEM-EDX) to probe the chemical distribution at the surface of the doped perovskite thin films on larger areas compared to the TEM specimens. We observe Na-rich regions that are mostly located at the grain boundaries across the entire sample area (Figure S12a and Figure S13a). For Cu- and Agdoped perovskites, we find a distribution of Cu and Ag across the entire area of perovskite films (Figure S13b, c and Figure S14b, c). We also detect regions rich in metallic Ag and Cu (Figure S13 b, c) that can be attributed to the higher concentration of dopants (e.g. 15 mol%) that were used in order to maximise the EDX signal being analysed.


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We performed periodic DFT calculations34 in atomistic-scale to examine how chemical bonding and band structure of the perovskite are affected by partial substitutions of a Pb with Na, Cu and Ag atoms (see SI for calculation details). The estimated formation energies for pristine and 3% Na-, Cu-, and Ag-perovskite single crystals (-30.95, -32.73, -41.25 and -66.86 kcal/mol, respectively) indicate preference for doping rather than phase separation. When the concentration of these cations is increased to 25%, as a result of the imbalance in chemical bonding due to the lower oxidation state of the monovalent cations, we simulate a preferential growth of the perovskite towards a 2D layered structure leading to a significant change in lattice and physical properties (See section A in SI). In the doped perovskites, we observe perturbations within the surrounding I6 octahedron due to the imbalance in chemical bonds i.e. different cation size and oxidation states of Na, Ag and Cu compared to Pb (Figure S15). Sodium is not able to bond to all the neighboring atoms (due to its ionic nature and smaller ionic radius), and we model it to form a single shorter spacing with one iodide atom, which tilts the I6 octahedron (Figure 6b, f). However, Cu and Ag can form 3 and 4 chemical bonds with iodide, respectively, due to their flexibility in rearranging themselves to the local environment with rearrangement of their d shell (Figure 6). Based on the distribution of the charge density in between the atoms, we can predict the nature of the chemical bonds. In the case of Na-I, bond charge densities are distinct and dense at the nuclei of Na and I, while for Cu-I and Ag-I, we observe a uniform distribution of charge densities between these atoms showing a covalent nature for these bonds (Figure 6e-h). Notably, an interplay of ionic radii and the electronegativity of the monovalent cations is responsible for the number and nature of the chemical bonds in between the surrounding iodide atoms and monovalent ions in the lattice.


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Figure 6. Schematic of the local lattice structure for the metal iodide octahedral (top panels) and the corresponding charge density plots (bottom panels) for the (a),(e) pristine, (b),(f) Na, (c),(g) Cu, and (d),(h) Ag-substitutionally doped CH3NH3PbI3 perovskites. Gray, purple, brown, light blue, cyan, orange and yellow spheres represent lead, iodide, carbon, nitrogen, sodium, copper and silver atoms, respectively.

In Table 2, we summarize the details of chemical bonds (bond length, tilting angle and average distortion factor) for pristine and doped perovskite structures. Since the chemical bonds between the dopants and iodide atoms are all in the ab-plane of the perovskite lattice, we calculate a small elongation of the lattice along the c-axis as well as an increase in the unit cell volume for Ag- and Cudoped perovskite (Table 2). We also find a different tilting for PbI6 octahedral in the vicinity of monovalent cations in the doped structures. This can slightly affect the electronic structure of the perovskite and the orientation of the MA dipoles.15 In Table 2, we also calculated the distortion factor (DF) defined as the averaged and squared difference of Pb−I bond lengths within the octahedron minus the average Pb−I bond length.46 It is clear that the DF increases upon incorporation of the dopants particularly for the Cu doped film. Page 14 of 32 ACS Paragon Plus Environment

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These increases originated from the broken symmetry in chemical bonding, which are reflected in the charge density plots near the dopants in Figure 6 e-h (e.g. the change in the bond length and number of attached iodide to each atom).

Table 2. DFT calculated average values of tilting angle, average distortion factor and bond length as well as the unit cell volume for pristine and 3% Na-, Cu- and Ag-substitutionally and interstitially doped perovskite.

Tilting Angle (Degrees)

Distortion Factor (a.u.)

Unit cell volume

Pristine

28.5

Substitutional

(Å )

Bond length (Å)

Horizontal bond length (Å)

Vertical bond length (Å)

1.74

991.8

3.21

3.22

3.19

24.5

2.21

991.8

3.20

3.21

3.19

Interstitial

27.6

1.53

976.9

3.20

3.21

3.20

Substitutional

26.0

3.50

1029.5

3.21

3.25

3.15

Interstitial

24.8

3.22

975.2

3.24

3.26

3.21

Substitutional

24.2

2.37

1028.7

3.24

3.22

3.27

Interstitial

22.5

1.85

972.5

3.22

3.26

3.18

Sample

3

Na

Cu

Ag

Furthermore, the simulated XRD spectra of the geometrically optimized pristine and doped perovskite single crystals (Figure S16), show a 0.5o shift of the main perovskite peak for the Ag and Cu-doped perovskite crystals, which is consistent with the GI-WAXS measurements. We also calculate the Partial Density of States (PDOS) of the doped and pristine perovskites and the corresponding band structures. We do not observe any contribution from the dopant orbitals near the band-edge features of the electronic structure, and conclude that doping does not introduce new energy states within the gap in agreement with the experimental data (Figure S17). The main contribution of dopant orbitals in the DOS is far away from the CB or VB edges. Therefore, the estimated band gap was the same for doped and pristine structures (1.58 eV). However, the dopants affect the curvature of the CB minimum and VB maximum (Figure S18) and thus the corresponding charge carrier effective masses, which are calculated using the standard effective mass equations and shown in Table S2 (see SI).


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Moreover, due to the high dipolar nature of the material with negatively charged octahedrons and positively charged MA cations together with different dipolar orientations of the MA cations inside the lattice, it is difficult to find a stable absolute reference for the vacuum level in the DFT calculations.34 Therefore, we estimate the amount of plausible VB and CB shifts based on the change in the VB edge with respect to the carbon core level. We then observe that the VB/CB edge of the Na, Cu- and Ag-doped perovskites shifted negatively with respect to the vacuum level (away from the vacuum level by 5 meV, 240 meV and 490 meV respectively) which is consistent with our XPS results (Figure 3e, f). The energy level of intra-band gap trap states, effective mass of charge carriers, position of Fermi level, nature of VB and CB edges and the structural parameters of the perovskite crystals greatly influence the photovoltaic performance of the subsequent devices.47 The driving potential for charge transfer from the perovskite layer to the charge selective contacts is directly proportional to the energy difference between the VB/CB edges of the perovskite and that of the electron/hole selective contacts. Therefore, optimization and further band alignment at the contacts is vital for obtaining high efficiencies for devices based on monovalent metal-doped-perovskite materials. Furthermore, our predicted Fermi level shifts show that Cu- and Ag-doped films have better charge balance in between the electron and hole densities and thus a superior charge collection efficiency compared to the pristine films. This could point towards efficient charge separation at the interfaces between perovskite and charge selective contacts. These are reflected directly in the enhanced optoelectronic and photovoltaic behaviour of the subsequent devices upon doping (Figure 1). The movement of the Fermi level towards the middle of the bandgap for the Cu- and Ag-doped films confirms dedoping of the highly n-doped pristine perovskite films, suggesting clearer bandgap consistent with the lower energetic disorder and sharper absorption edge for the doped perovskite films (Figure S1). Since exact location of the dopants within the perovskite lattice in atomistic scale cannot be deduced from our experimental results, the incorporation of the monovalent cations at interstitial sites would not be excluded. Therefore, we investigate the possibility of the interstitial doping via allocation of the abovementioned monovalent cations in a preferential location and relaxing the position of all the Page 16 of 32 ACS Paragon Plus Environment

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atoms inside the unit cell accordingly (Figure S19). In Figure 7, we show formation of chemical bonds between monovalent cations and iodine atoms, where Cu and Ag atoms are located within the same plane of Pb and equilateral iodine atoms while Na shows a displacement out of this plane. The formation energies for interstitial doping (e.g. 3%) of Na, Cu and Ag are -38.25, -70.43 and -75.82 kcal/mol, respectively, which are slightly more favorable compared to substitutional doping indicating the possibility of coexistence of both type of doping in the perovskite lattice. In Figure 7, we also show the charge density plots for the interstitial doping, indicating imbalanced charge distribution near to the covalent bonding of monovalent cation dopants and iodide. We find no intra band gap states for interstitial doping and a slight change in effective mass of charge carriers with no impact on the optical band gap similar to the substitutional doping (Figure S20). We also summarise the chemical bond details upon interstitial doping in Table 2.

Figure 7. Schematic of the local lattice structure for the metal iodide octahedral (top panels) and the corresponding charge density plots (bottom panels) for the interstitial doping of CH3NH3PbI3 perovskite with (a),(d) Na, (b),(e) Cu, and (c),(f) Ag. Gray, purple, brown, light blue, cyan, orange and yellow spheres represent lead, iodide, carbon, nitrogen, sodium, copper and silver atoms, respectively. The white arrows indicate the charge imbalance close to Pb-I bonds. Page 17 of 32 ACS Paragon Plus Environment

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In brief, incorporation of the abovementioned monovalent cations has shown interesting impacts on the film growth, chemical and electronic structure of the perovskite confirming high tolerance of this material to additives. In the case of Na doping similar to other alkali metals such as potassium23 or rubidium,48 the monovalent cation decorates the grain boundaries and surfaces passivating the surface defects with minimal effect on the band structure of the perovskite. However, for Cu and Ag cations, the more uniform distribution and incorporation into the perovskite lattice either in the interstitial or substitutional (e.g. with Pb2+) positions is plausible with significant impact on the electronic structure (e.g. dedoping). Expectedly, these type of doping need to be accompanied by relevant charge compensation processes such as interstitial occupation of iodide or a combination of both of the doping processes. Notably, substitutional doping can compensate the halide vacancies leading to the observed shift of the Fermi level. However, the incorporated amount are far larger than required to compensate the iodide vacancies. Therefore, they must be incorporated in such way that the net charge become zero. This can be achieved via involving the presence of further halide vacancies in the interstitial sites or/and balance of interstitial and substitutional with Ag or Cu doping. CONCLUSIONS In summary, we have demonstrated significant enhancement in optoelectronic and photovoltaic properties of perovskite semiconductors doped with monovalent cation halide (e.g. Na, Cu and Ag). We have investigated the influence of these cations on the growth, and crystal structure of lead halide perovskite with a comprehensive insight on the chemical distribution of these dopants within the lattice and their effect on the band structure. We observe a distinct shift in the main diffraction peak of both PbI2 and CH3NH3PbI3 upon addition of the above-mentioned monovalent cations, which is a sign of partial incorporation in the perovskite structure. We then detected all the dopants via soft XPS and HAXPES and we recognize that the concentration of Cu is significantly higher in the bulk compared to the surface in contrast to the other dopants. We detect the highest inclusion for Na with the least active interaction into the perovskite band structure as per the overlap between their valence band edges with the pristine perovskite films. However, we show that introducing Cu and Ag into the Page 18 of 32 ACS Paragon Plus Environment

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perovskite film de-dope the highly n-doped pristine perovskite by moving the Fermi level towards the middle of the gap leading to cleaner bandgaps as reflected in the lower energetic disorder for the doped perovskite films. From the STEM-EDX results, we find formation of Na-rich phases at the grain boundaries and top/bottom interfaces while Ag-doped perovskite showed Ag-rich inclusions within the perovskite films. Our DFT calculations show how the chemical bonds, lattice arrangement and band structure of perovskite will respond to the substitutional and interstitial doping of perovskite with a monovalent cation (e.g. Na, Cu and Ag) in atomic scale. We also find that interstitial doping is more favorable than substitutional doping based on the calculated formation energies despite that a combination of both type of doping is more probable to happen. Our findings show a fundamental insight for the local chemistry and changes due to the monovalent metal doping and the corresponding impacts on the electronic structure that guide the researchers for further in-depth investigations. METHODS Materials. All the organic cation salts were purchased form Dyesol. The Pb compounds were purchased from TCI. Unless otherwise stated, all other materials were purchased from Sigma-Aldrich. Perovskite film preparation and device fabrication. PbI2 was dissolved in DMF at a concentration of 1 mol/L under constant stirring at 80°C. The halide salts including NaI, CuI, and AgI were then added to the lead halide solution at a concentration of 0.03-0.08 mol/L. All the solution were prepared inside an argon glove box under moisture- and oxygen-controlled conditions. The CH3NH3PbI3 films were prepared using a two-step deposition method.49 In the first step, the pure or additive based PbI2 solution was spin-coated at 6500 rpm for 30 s and dried at 80°C for 30 min. To form the perovskite, a solution of CH3NH3I in 2-propanol (8 mg/mL) was dropped with a pipette on top of the PbI2 coated substrate mounted on a spincoater, left for 30 s, and then spin-cast at 4000 rpm for 20 seconds. The resulting films were dried at 80°C for 30 minutes. We prepared all solutions and films in a nitrogen-filled glove box. For solar cell fabrication, we followed the same procedures for substrate preparation as well as deposition of both electron and hole transport layers (i.e. TiO2, SpiroOMeTAD with additives) as in our earlier works.49


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Photothermal deflection optical absorption spectroscopy (PDS). PDS is an optical absorption technique, which measures absorption directly by probing the heating effect in the samples upon absorption of light resulting in its capability to measure an absorption, which is 5 - 6 orders of magnitude weaker than the band-edge absorption. The samples used for the PDS measurements were prepared in an identical fashion to the solar cell preparation and spun onto water free quartz substrates (which were cleaned with acetone, isopropanol, and deionized water followed by a 10 min oxygen plasma etch). For this particular measurement, we made use of quartz rather than the FTO-coated glass to minimize light absorption because of the substrate. During the measurement, we kept the samples in a hermetically sealed quartz cuvette filled with an inert liquid, Fluorinert FC-72 from 3M Corporation, which acts as the deflection medium with high temperature dependent refractive index. We excited the perovskite films with a modulated monochromated light beam perpendicular to the plane of the sample. A modulated monochromated light beam was produced by a combination of a Light Support MKII 100W Xenon arc source and a CVI DK240 monochromator. The transverse probe beam was produced with Qioptiq 670-nm fibre-coupled diode laser and passed as close as possible to the sample surface. Beam deflection was measured using a differentially amplified quadrant photodiode and a Stanford Research SR830 lock-in amplifier, which is proportional to the absorption in the sample. Time-resolved photoluminescence. Time-resolved photoluminescence measurements were acquired with a gated intensified CCD camera system (Andor iStar DH740 CCI-010) connected to a grating spectrometer (Andor SR303i). Excitation was performed with femtosecond laser pulses which were generated in a homebuilt set-up by second harmonic generation in a β-barium borate crystal from the fundamental output (pulse energy 1.55 eV, pulse length 80 fs) of a Ti:sapphire laser system (Spectra-Physics Solstice). Temporal resolution of the photoluminescence emission was obtained by measuring the photoluminescence from the sample by stepping the intensified CCD gate delay relative to the pump pulse. The gate width was 20 ns. Photoluminescence characterization. Photoluminescence quantum efficiency (PLQE) measurements were taken by mounting perovskite films or encapsulated device stacks in an


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integrating sphere and photoexciting with a 532-nm continuous-wave laser at the given intensity. The laser and the emission signals were measured and quantified using a calibrated Andor iDus DU490A InGaAs detector for the determination of PL quantum efficiency. The external PLQE was calculated as per de Mello et al.50 Solar cell characterisation. Current–voltage characteristics were recorded by applying an external potential bias to the cell while recording the generated photocurrent with a digital source meter (Keithley Model 2400). The light source was a 450 - W xenon lamp (Oriel) equipped with a Schott-K113 Tempax sunlight filter (Praezisions Glas & OptikGmbH) to

match the emission

spectrum of the lamp to the AM1.5G standard. Before each measurement, the exact light intensity was determined using a calibrated Si reference diode equipped with an infrared cut-off filter (KG-3, Schott). All measurements were conducted using a non-reflective metal aperture of 0.105 cm2 to define the active area of the device. Conductivity measurements. To measure conductivity, space charge limited current (SCLC) in a vertical configuration was adopted. Perovskite films (~400 nm) were deposited on 25 nm gold coated corning 1737F glass substrate (15 mm × 15 mm). A top electrode of 25 nm gold was deposited on top of the film. The films were characterized using Keithley 4200 SCS at room temperature. Grazing Incidence Wide Angle X-Ray Scattering (GIWAXS). Part of the GI-WAXS measurements were performed at the XMaS beamline (ESRF, France). A fixed-exit, water-cooled, double crystal Si(111) monochromator, placed at 25 m from the source was used to monochromatize the X-ray beam coming from a bending magnet (Ec = 9.8 keV). The X-ray energy was tuned to 10 keV (1.2398 Å) and a Rh-coated toroidal mirror was used to focus the monochromatic beam horizontally and vertically. The beam flux was ~5×1010 photons/s at the sample position. The original beam spot size was 500 (horizontal) × 400 (vertical) µm2 at the sample position. We employed a set of motorized slits (Huber, Germany) immediately before the sample to have a better-defined footprint in the vertical direction. The final beam spot size with slits was 300 (horizontal) × 115 (vertical) µm2. The beam footprint extended 300 µm horizontally and throughout the perovskite films. The samples


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were measured at an out-of-plane incident angle of ~0.3°, corresponding to a penetration depth of ~140 nm, probing a significant portion of the bulk. The other part of the GI-WAXS measurements were performed with a XEUSS (XENOCS, France) laboratory beamline located at the SMALL Lab at the University of Sheffield. We employed a liquid gallium source (9.2426 KeV). Diffraction patterns were collected with a Pilatus3 1M detector (Dectris) at an ouf-of-plane angle of ~0.3°. Photoelectron Spectroscopy. In-house X-ray Photoelectron Spectroscopy (XPS) and Hard X-ray Photoelectron Spectroscopy (HAXPES) were performed in this work.

In-house XPS

measurements were carried out with a Scienta ESCA 300 instrument, using monochromatized AlKα radiation (hν=1486.6 eV). HAXPES was performed at BESSY II (Helmholtz Zentrum Berlin, Germany) at the KMC-1 beamline51 using the HIKE end-station.52 To obtain a deeper depth level than that of classical XPS, an incident photon energy of 4000 eV was used by selecting the first-order light from a Si(311) crystal of a double-crystal monochromator (Oxford-Danfysik). Calibration was done by measuring a gold plate reference and setting the Au 4f7/2 core level peak to 84.0 eV, corresponding to a zero binding energy at the Fermi level. Quantification and intensity ratios were calculated from the experimental results after correcting the intensity by the photoionization cross section for each element at their specific photon energy, using Scofield’s relative sensity factors database values. The samples investigated by XPS and HAXPES were the same and stored in a sealed box with common desiccants to limit moisture contamination. Electron Microscopy. Samples for STEM-EDX were prepared as lamellae using a focused ion beam (FEI Helios Nanolab). HAADF images and elemental maps were acquired in a FEI Tecnai Osiris S/TEM, operated at 200 kV. Each map had a pixel size of 10x10 nm2. The maps were analysed using Hyperspy, a python-based open source toolkit for electron microscopy data analysis. Denoising was carried out using Principal Component Analysis (PCA) within Hyperspy. Scanning electron microscopy – EDX was carried out on a FEI Nova NanoSEM using a Bruker XFlash 6 detector. The maps were processed within Esprit 2.1.


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DFT Calculations. The crystal DFT calculations were performed following the previously published procedure.34 The generalized gradient approximation (GGA) for exchange correlation using the Perdew-Burke-Ernzerhof (PBE) functional was applied and simulations used pseudopotential53 as implemented within the Quantum Espresso package.54 The Silver 5d10/6s1,Copper

4d10/5s1,

Sodium 2p6/3s1, lead 5d10/6s2/6p2, the nitrogen 2s2/2p3, the iodide 5s2/5p5, and the carbon 2s2/2p2 electrons were considered as valence electrons. The super cells consisted of 384 atoms in a 2x2x2 tetragonal Bravais lattice, and the initial cell parameters were set to a = 16.42 °A and c = 24.92A° for CH3NH3PbI3. The tetragonal unit cell can be described as composed of four slightly distorted cubic unit cells with a cell parameter of around 6.31 A°. Substitution of one lead with Ag, Cu or Na within the lattice corresponds to the ̴ 3% doping. For the interstitial doping, the unit cell has 385 atoms and the dopant has been placed close to an estimated energetically favorable cite within the lattice. The unit cell volume and all coordinates of the atoms have been optimized to have a total force lower than 0.04 Ry/Boh. Cutoff energies for the plane wave function and the charge density were set to 40 and 450 Rydberg, respectively. Brillouin zone samplings were carried out by a 2×2×2 Monkhorst-Pack grid. Full details of the DFT calculations and experimental methods can be found in the Supporting Information. ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publication website at DOI:XXX. The PDS absorption spectra of doped perovskite films using two different deposition methods, current density-voltage characteristics and the corresponding photovoltaic parameters of champion dopedperovskite solar cells, GI-WAXS diffraction patterns, top view SEM images for doped perovskite and lead-iodide, the HAXPES overview spectra, high resolution core level spectra and valance band spectra of pristine and doped perovskite thin films, conductivity measurements for the doped perovskite films, the HAADF STEM cross sectional images and the corresponding NMF decomposition results for the doped films, SEM images and SEM-EDX elemental maps of doped


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perovskite films, details for DFT calculations and the resulting calculated structures, simulated XRD spectra, PDOS, band structure for 3% doping, DFT calculation results for 25% doping of perovskite with monovalent metals are included in the supporting information.

ACKNOWLEDGMENTS M.A.-J. thanks Nava Technology Limited and Cambridge Materials Limited (CML18030501) for their funding and technical support. The authors thank the Engineering and Physical Sciences

Research Council (EPSRC) for financial support (grant number: EP/M005143/1). B.P. and H.R. acknowledge financial support from the Swedish Research Council, the Swedish Energy Agency, and the Swedish Foundation for Strategic Research as well as StandUP for Energy. HZB is acknowledged for the allocation of synchrotron radiation beamtime. M. A. thanks the president of the UAE's Distinguished Student Scholarship Program (DSS), granted by the Ministry of Presidential Affairs. M.I.D, and M.G. thank the King Abdulaziz City for Science and Technology (KACST), and acknowledge funding from the European Union's Horizon 2020 programme, through a FET Open research and innovation action under grant agreement number 687008. XMaS is a mid-range facility

supported by EPSRC. We are grateful to all the XMaS beamline team staff for their support. We thank Dr. Andrew Parnell and Dr. Oleksandr Mykhaylyk for their support with the XEUSS system at the SMALL Lab (University of Sheffield). We thank Prof. David Lidzey for providing financial support for accessing the SMALL Lab (EPSRC grant number: EP/M025020/1). DFT calculations in this work made use of computational resources provided by the Swedish National Infrastructure for Computing (SNIC) through Uppsala Multidisciplinary Center for Advanced Computational Science (UPPMAX) under Projects snice 2017-01-15 and snice 2016-10-23. MP appreciates discussions and support from Prof. Kersti Hermansson, Peter Broqvist and Prof. Tomas Edvinsson.


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