The surface of a semiconductor is where the periodic structure of the crystal lattice is terminated, accompa-nied by reconstruction of the atomic structure, owing to surface- energy minimization. Interruption of the periodicity, either intrinsically or extrinsically, gener-ates surface- specific states that deviate from those in the bulk crystal1,2. These states substantially affect the performance of semiconductor microelectronic devices, motivating intensive research towards understanding the surface science of semiconductors. A new class of semiconductors, metal halide perovskites (MHPs), has emerged in the past decade to prevail in the optoelec-tronics landscape3–5. Exploiting the exceptional optical and electronic properties of these materials has led to the remarkably rapid development of MHP- based devices. MHP- based photovoltaic cells with power- conversion efficiencies (PCEs) of >25% and light- emitting diodes (LEDs) with external quantum efficiencies of >20% were developed within a decade6–9. The general chemical for-mula of MHPs is ABX3, where A+ is a monovalent cat-ion, such as methylammonium (MA+), formamidinium (FA+) or caesium (Cs+); B2+ is a divalent cation, such as Pb2+ or Sn2+; and X− is a halide anion. Compared with classical perovskite oxides, the larger and more polar-izable ionic components, as well as the lower charge on the X- site, result in weaker coulombic interactions and softer ionic bonding within MHPs. These softer interac-tions produce a mobile and sensitive surface that facil-itates atomic reconstruction and irregularities. Various combinations of ions can occupy the three sites, gov-erned by the geometrical Goldschmidt tolerance factor

(for close- packed ions), giving rise to compositional tun-ability. Mixed- ion occupation is allowed for each site, further extending the tunability of the chemical compo-sition. The structural susceptibility and compositional variability give rise to the intricate surface behaviour of MHPs, making the role of the surface more pronounced in affecting the materials properties and device perfor-mance. Lattice discontinuities and perturbation of the periodicity of chemical bonds, intrinsically or by exter-nal stimuli, at the surface of MHPs creates electronic states specific to the surface that can have a substantial effect on the band structure and electron behaviour. Moreover, the structural disorder at the surface is gen-erally metastable, intensifying degradation pathways of the materials and devices10,11.

As the surface- to- volume ratio scales inversely with the linear dimensions, surface effects intensify as the particle size shrinks and eventually become dominant in quantum dots (QDs), which are a few nanometres in size. The large contribution from the surface means that the properties of semiconductor QDs are significantly influenced by the surface features12. Nevertheless, their unique ‘defect tolerance’ (discussed below) enables MHP QDs to be highly emissive, even without a second surface- capping material to form the core–shell structure that is widely adopted in conventional QDs. However, most of the photoluminescence quantum yields (PLQYs) reported are only ‘close’ to unity, indicating that MHPs are defect tolerant rather than entirely defect free13. This finding further emphasizes the significance of the sur-face in enhancing the optical properties of MHP QDs,

The surface of halide perovskites from nano to bulkJingjing Xue , Rui Wang and Yang Yang  ✉

Abstract | The surface of a semiconductor often has a key role in determining its properties. For metal halide perovskites, understanding the surface features and their impact on the materials and devices is becoming increasingly important. At length scales down to the nano scale regime, surface features become dominant in regulating the properties of perovskite materials, owing to the high surface- to- volume ratio. For perovskite bulk films in the micrometre range, defects and structural disorder readily form at the surface and affect device performance. Through concerted efforts to optimize processing techniques, high- quality perovskite thin films can now be fabricated with monolayer- like polycrystalline grains or even single crystals. Surface defects therefore remain the major obstacle to progress, pushing surface studies to the forefront of perovskite research. In this Review, we summarize and assess recent advances in the understanding of perovskite surfaces and surface strategies towards improving perovskite materials and the efficiency and stability of perovskite devices.

Department of Materials Science and Engineering, University of California, Los Angeles, Los Angeles, CA, USA.

✉e- mail: yangy@ucla.edu

https://doi.org/10.1038/ s41578-020-0221-1

REVIEWS

Nature reviews | Materials

especially in optoelectronic applications for which the surface- ligand chemistry needs to be adjusted to provide a benign environment for charge- carrier transport14. As the length scale increases to micrometres, the regime reaches that of bulk films, upon which the most effi-cient MHP photovoltaic devices to date are based. The unremitting efforts to establish sophisticated processing techniques have culminated in high- quality MHP thin films with monolayered polycrystalline grains or even

single crystals15–17. Considering the vertical direction of carrier transport in photovoltaic and LED devices, the surface imperfections are now the primary obstacle in limiting the carrier dynamics and, thus, device perfor-mance. Most of the recently reported high- efficiency perovskite optoelectronic devices have been enabled by surface- treatment techniques7,9,15,17–19, signifying the crucial role of surfaces in the most advanced perovskite devices.

Taking advantage of the defect tolerance of MHPs, high- performance MHP devices have been achieved at an early stage; however, the impossibility of defect- free MHPs, according to the maximum- entropy principle20, has brought the study of MHP surfaces to the fore-front. In this Review, we highlight the role of surfaces in MHPs, starting with a discussion of the fundamental understanding of the molecular, atomic and electronic structures at the surface. The physicochemical picture of MHP surfaces is constructed by examining MHPs of varying size, ranging from QDs to bulk films. We then assess the state- of- the- art techniques and approaches used to characterize surface features, before providing an overview of the advances towards higher- performance MHP- based optoelectronic devices. Finally, we consider the challenges in obtaining in- depth understanding of MHP surfaces and deliberate possible approaches to overcoming these issues and further enhancing the performance of MHP optoelectronic devices.

Surface structures and propertiesSurface terminations and disorderThe 3D crystal structure of MHPs consists of corner- sharing [BX6]2− octahedra with an A- site cation occu-pying the central void. When the atomic periodicity is interrupted, the crystal surface can terminate with either metal halide units or organic cations. As the majority of surfaces are metastable and MHPs possess a particu-larly weak ionic bonding character, the surface termina-tions of MHPs are highly dependent on their synthesis and/or processing history. Different surface terminations are associated with different surface properties, which, in turn, impart the character of the entire material and the device behaviour. Hence, for the development of more efficient devices, it is prerequisite to have a comprehen-sive understanding of the atomic structure of the surface terminations and to establish surface structure–property relationships.

Surface energetics. In terms of lowest- energy config-urations, the (001) and (110) surfaces of MAPbI3 are generally considered to be energetically favourable21,22. Three representative surface terminations on a (001) surface are MAI, PbI2 flat and vacant21 (Fig. 1a). Focusing on the (001) surface, PbI2 flat termination is less thermo-dynamically stable than MAI termination23. Theoretical analysis of the surface grand potentials of various sur-face terminations of PbIx polyhedra found that vacant and PbI2 flat terminations coexist on the energetically favourable (110) and (001) surfaces24. Nevertheless, under the thermodynamic equilibrium conditions of bulk MAPbI3, a vacant termination is generally more stable than PbI2 flat terminations on all surfaces24,25.

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Fig. 1 | surface terminations and disorder. a | Side views of three representative surface terminations in MAPbI3 (where MA+ is methylammonium). b | The 12 types of native point defects found in metal halide perovskites. VY denotes a Y vacancy, Yi denotes an interstitial Y- site and YZ denotes a Z- site substituted by Y, where Y and Z represent an ion of APbX3 (where A+ is a monovalent cation and X− a halide). c | Other surface features of metal halide perovskites include ion dimers (Pb2+ dimers and halide dimers), zigzag patterns, various orientations of the organic cations and the incursion of foreign species, such as O2. The high- resolution scanning tunnelling microscopy images on the right show the formation of halide dimers (top) and a zigzag pattern (bottom) on the surface of MAPbI3. The unit cells are denoted by dashed white rectangles. Panel a adapted with permission from reF.21, ACS, further permissions relating to the material excerpted should be directed to the ACS. Panel c (scanning tunnelling microscopy images) adapted from reF.36, ACS.

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Similar thermodynamic energetics were observed for the MABr- terminated (001) surface of MAPbBr3 (reFs26,27).

Considering the two dipole orientations of polar ammonium cations at a surface, a lower surface energy is achieved when the dipoles point into the inorganic layer26. By starting from a cubic geometry and allow-ing the MAPbI3 bulk cell to relax, insight into the (001) surface of a pseudocubic structure was obtained using molecular dynamics simulations. The surface relaxation was driven by competition between contraction of the inorganic cage on the PbI2 flat (001) surface, maximi-zation of the hydrogen bonding between MA+ and the bridging I− ions, and the dipole alignment of the MA+ cations28. Owing to the weak binding of the ionic lattice and the dynamic rotation of the organic cations, the sur-face terminations of MHPs do not have a fixed composi-tion or structure, particularly non- stoichiometric MHP surfaces with various types of structural disorder, which trigger more intricate surface reconstruction.

Intrinsic surface disorder. As in the bulk, there are three types of intrinsic point defects that are likely to exist at the surface of MHPs, namely, interstitial defects, antisite defects and vacancies29–31. Overall, 12 point defects can be present at a MHP surface (Fig. 1b). Theoretical analy-sis of the unusual defect physics in MAPbI3 suggests that the dominant point defects create only shallow trap states, whereas the remaining point defects that are likely to create deep- trap states have high formation energies, which is regarded as the origin of the defect tolerance in MHPs30,32. However, when these defects are present on the surface, their energetics can become more com-plicated and might deviate from those in the bulk. The chemical environment around a defect on the surface differs from that in the bulk, owing to variations in the surface terminations, the chemical conditions and the undercoordinated sites with reduced steric hin-drance. The relationship between surface- defect for-mation and the surface chemical conditions has been investigated in MAPbI3 (reF.20). It was suggested that different types of defects are dominant on the surface, depending on the chemical conditions, that is, differ-ent surface terminations form under I- rich, Pb- rich or moderate (thermodynamic equilibrium with the bulk MAPbI3) conditions. The formation energies of carrier- trapping surface defects are generally higher under Pb- rich conditions than under I- rich conditions, indicat-ing the superiority of the Pb- rich over the I- rich environ-ment. Under the moderate condition, all surface point defects except Pb vacancies, which are shallow carrier- trapping states, show high formation energies, suggest-ing that this condition might be advantageous to reduce the number of detrimental trap states on the surface21.

Interest in MHPs based on FA+ is growing, owing to the possibility of achieving higher PCEs. Theoretical modelling indicates that, compared with MA+- related defects in bulk MAPbI3, FA+- related defects in FAPbI3 form more readily, owing to weaker van der Waals inter-actions between the FA+ cations and [PbI6]2− octahedra33. This weaker binding results from the larger size and smaller molecular dipole moment of FA+ compared with MA+ (reF.34). The defect formation energies in the

bulk were further compared with those on the surface of FA+- based perovskites to investigate the dominant sur-face defects. With a much lower formation energy, the antisite PbI defect was found to more readily form on the surface than in the bulk15.

The calculation of the transition levels of charged defects on surfaces is, however, complicated by the dis-sipation of charges, which limits the accuracy of theo-retical predictions of MHP surface defects and increases the significance of experimental approaches to probe the atomic structure of perovskite surfaces. Distinct surface features, such as halide dimers (accompanied by unpaired halide protrusions) and zigzag patterns were experimen-tally observed on the surface of MAPbX3 using scanning tunnelling microscopy (STM; Fig. 1c, top right). These features most likely originate from the reorientation of the surface MA+ cations and the electrostatic interaction between the positive charge of the ammonium groups and the neighbouring halides35–37. Two different sur-face structures with stripes and armchair domains were recently observed in all- inorganic CsPbBr3, originating from a complex interplay between the Cs+ cations and Br− anions38. Vacant defect clusters were also detected in MAPbBr3, and identified as dynamic pairs of MA+ and halide vacancies that tend to diffuse together when moving along a specific crystallographic orientation. Ion transport on or towards the surface was assisted by these vacancy clusters37, indicating that the surface disorder of MHPs is far from simple combinations of intrinsic point defects or defect pairs. Intricate defect coupling, grouping and/or reconstructions are likely to be widely present given the flexibility of the surface, with more space for geometrical reorganization than that in the bulk37,39,40. Structural heterogeneity observed in bulk MHPs, such as twinning, Ruddlesden–Popper planar faults and self- organized superlattices with phase coexist-ence, would also affect surface structures41–43. Therefore, further studies towards a more comprehensive picture of surface states are in high demand.

Extrinsic surface disorder. The metastable nature of MHP surfaces makes them susceptible to attack by foreign species, resulting in various types of extrinsic surface disorder, which contribute to the complexity of perovskite surface structures (Fig. 1c, bottom left). For example, under ambient atmosphere, the surfaces of MHPs suffer from the intrusion of oxygen and moisture. The photoinduced oxygen- degradation mechanism of MAPbI3 involves the simultaneous reduction of O2 in I− vacancies and the formation of a superoxide44. Various surface features are observed upon exposure of a MHP to moisture, owing to the strong surface- termination dependence of the interaction with water molecules. For example, the MAI- terminated surface of MAPbI3 undergoes rapid hydration driven by the absorption of water molecules at the Pb sites. By contrast, the PbI2- terminated surface was more immune to surface sol-vation by virtue of the moderately robust Pb–I bonds, except when the termination became defective and facilitated the hydration process45,46. Water absorption on MAPbI3 surfaces also shows polarization depend-ence, with absorption being more favourable on an

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MAI- terminated surface with −NH3+ pointing towards

the surface than on an MAI- terminated surface with −CH3 pointing towards the surface. By contrast, the PbI2- terminated surface shows the opposite dependence on the dipole orientation47. In summary, although the MAI- terminated surface is thermodynamically more stable, the MAPbI3 surface tends to reconstruct into a PbI2- terminated one, ultimately generating a hydrated species upon exposure to ambient conditions.

Electronic structure and propertiesThe atomic structure of the surface determines the dis-tribution of surface electronic states. Theoretical simu-lations reveal that the electronic properties at the surface of MHPs do not significantly deviate from those of the bulk. Particularly for non- defective or vacancy- only surface terminations, the surface electronic structures largely maintain the characteristics of the bulk phase, with no midgap states observed24–26,31. The band structure at the surface of MAPbI3 resembles that of the bulk form, with the valence band maximum consisting mostly of I 5p and Pb 6s orbitals, and the conduction band mini-mum largely derived from Pb 6p orbitals24,48,49. Owing to the absence of strong interactions between the MA+ cations and [PbI6]2− network, the hybridized σ bonds of MA+ are located deep within the valence band and remain without major dispersion. The strong coupling between the lone- pair Pb 6s and the I 5p orbitals raises the valence band maximum and, hence, acceptor- like defects generally end up shallower than the band edge. The high ionicity also makes the energetically dominant donor- like defects shallow- level states31. Consequently, defects with a low formation energy create only shallow trap states; this is interpreted as the primary origin of the so- called defect tolerance30,48,50,51. However, when look-ing into the extraordinarily long carrier lifetime within a MHP single crystal52, compared with those in a poly-crystalline film, the question arises as to what extent the surface electronic properties resemble the bulk properties and whether the MHP surface is indeed defect tolerant.

Disregarding the effects from surface lattice irregu-larities, termination- dependent variations in the band structure have been observed. In MAPbX3, the band-gap of the PbX2- terminated surface is slightly smaller than that of the MAX- terminated surface, owing to the emergence of surface states located just above the valence band maximum24–26,53. The surface- weighted states do not fall in the midgap of the band structure, and, thus, they can probably be characterized as sur-face resonant states rather than true surface states and could be beneficial to hole transfer24–26. For MAPbBr3, surface reconstruction and the interplay between the orientation of the organic cations and the positions of the hosting halide anions was found to lead to the forma-tion of electronically distinct domains with ferro electric and antiferroelectric character, and contributions from both Br− and MA+ orbitals to the electron density on the surface35. Additionally, the Rashba effect — that is, momentum- dependent band splitting as a conse-quence of spin–orbit coupling and inversion symmetry breaking54 — also exhibits strong surface dependence in MHPs. In contrast to the ‘dynamic’ Rashba effect, with

fluctuations on a temporal and spatial scale, reported for the bulk of MHPs55, a ‘static’ band- splitting effect is found on the surface and is probably due to sur-face structural distortions56,57. The reorientation of the organic cations in the outermost layers (Fig. 1c, bottom centre) in MAI- terminated surfaces only marginally affects the inorganic cages but causes pyramidalization of the outermost Pb cations in PbI2- terminated surfaces, which increases the Rashba splitting56. The effect of MA+ alignment in the bulk, however, leads to stronger Rashba splitting than that on the surface. For mixed- halide MA(Br1−xClx)3 (0 ≤ x ≤ 0.19) single crystals, the Rashba effect was suppressed at the surface, owing to the dis-order in the MA+ alignment, but increased in the bulk upon Cl− incorporation, as a result of compression of the unit cell volume and, thus, the space available for MA+ orientational degrees of freedom58.

Surface structural disorder introduces specific fea-tures in the electronic energetic landscape. As evidenced using transient reflection spectroscopy, the total carrier lifetime in a polycrystalline MHP thin film is limited by the recombination at the top and bottom surfaces, indi-cating that many detrimental trap states are present at the surfaces, despite the defect tolerance observed in the bulk59. As the defect formation energy is highly depend-ent on the chemical environment21,60, it is plausible that the proposed deep- level point defects of the bulk with high formation energies could form more readily on the surface, owing to the altered chemical environment, engendering many non- radiative charge- recombination centres that are particular to the surface21,40,61. The ener-getically favourable defects of the bulk that do not affect the band structure, such as vacancies, cluster on the sur-face and considerably raise the local work function of the perovskite surface. Additional surface electronic fea-tures could develop through, for example, the covalency- induced formation of ion dimers62 (Fig. 1c, top left) and vacancy- based doubly charged surface layers63. The invasion of foreign species under ambient conditions also contributes to the distortion of the electronic structure at the surface. The incorporation of oxygen into defect- free MAPbI3 creates midgap trap states that originate from the unoccupied oxygen π antibonding orbital, and the occupation of I− vacancies with O2

− superoxide species shifts the oxygen state downwards into the valance band, indicating favourable oxygen reduction44. The absorp-tion of water molecules on the surface slightly increases the surface bandgap and induces more degrees of disor-der in the surface structure of MHPs. External stimuli, such as incident light, also induce specific surface states in MHP thin films. Irradiation of (Cs0.06MA0.15FA0.79)Pb(Br0.4I0.6)3 films triggered the formation of inhomo-geneous I−- rich perovskite domains at the surface, leading to low- bandgap emitting sites with a high local concentration of charges, ensuring radiative bimolecular processes dominate recombination64.

Surface features of MHP QDsThe surfaces of QDs are distinct from those of bulk counterparts because the ligands used in the synthe-sis contribute substantially to the surface structure. According to the covalent- bond- classification method,

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the surface ligands of QDs include L- type (two- electron donor ligands, which act as Lewis bases), Z- type (two- electron acceptor ligands, which act as Lewis acids) and X- type (one- electron donor ligands)12,65 (Fig. 2a). Most of the MHP QDs synthesized to date primarily rely on long- alkyl- chain ligands, the prototypical example of which is the oleylamine–oleic acid pair66. With the ionic character of MHP QDs, the surface ligand binding falls into the binary X- type category; surfaces are typically

dynamically stabilized with either an oleylammonium halide or oleylammonium oleate67. The cation and anion bind onto the surface of MHP QDs as an ionic pair, with the ammonium cation taking the place of a surface A- cation, while the halide or carboxylate attaches on the surface and maintains charge neutrality67–69 (the blue structures in Fig. 2b). The large diffusion coeffi-cient of surface ligands implies highly dynamic ligand binding, with fast exchange between the bound and

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free states, which is likely the origin of the ease of halide exchange39,70 and ligand loss upon purification of MHP QDs71–73. This unique binary binding mode has inspired numerous new ligands for MHP QDs (Fig. 2b), the details of which are discussed below.

The non- unity and frequent loss of PLQY upon cyclic purification implies the existence of structural disorder at the surface of MHP QDs. Surface imperfections are more prominent in nanoparticles than in the bulk phase, owing to their larger number of surface atoms and size of a few nanometres. The small size enables bulk defects in the core of the QDs to travel towards the surface if they are more stable in this region. The energetics of surface- defect formation in QDs are also expected to differ from those in the bulk, as the chemical environ-ment of QD surface defects involves a complex inter-play between lattice termination, ligands, surrounding solvent and size- induced inherent strain redistribution and symmetry breaking74,75. The most plausible scenario for the formation of defects on the surface is the detach-ment of capping ligands to leave behind vacancies. The electronic structure evolution of [CsPbX3](PbX2)k{AX′}n (where k and n are the numbers of PbX2 units and {AX′} ligands on the two outermost atomic layers, respec-tively) was evaluated theoretically while stripping ion pairs from the QD surface. Even when 75% of the {AX′} units had been removed, midgap states remained absent and the structural integrity was preserved. However, such defect- tolerant behaviour no longer holds for PbX2- terminated [CsPbX3](PbX2)k or uncapped QDs, for which stripping only 25% of the surface PbX2 units causes midgap states to appear and triggers structural deformation76. Deeper insight into the electronic struc-ture of surface defects in MHP QDs was recently gained through systematic theoretical screening of interstitial, vacancy and antisite defects on the surface of CsPbBr3 QDs75, as has been intensively studied in the bulk. In con-trast to bulk MHPs, nanoparticles capped with surface ligands are usually formed in a low dielectric solvent. Hence, ion pairs are usually responsible for the creation of the point defects, where the addition, removal or sub-stitution of an ion on the surface has to be accompanied by a counterion to neutralize the charge. Therefore, the

atomic and electronic surface features generated in nan-oparticles can be distinct from those of the bulk. The defect formation energies of each type of point defect at different locations were compared (Fig. 2c). With the most negative defect formation energies, surface intersti-tial defects were determined to be energetically favoured, especially at the surface edge. Midgap states of the most stable Pb2+ interstitials emerged from the counterion (halide) instead of Pb2+ itself, and only when the halide occupies a surface position that points outwards from the surface. Although the expulsion of ions in their elemental form from the lattice would also result in midgap states, this process is strongly endothermic, which limits the formation of such surface defects. Other point defects on QD surfaces either do not create midgap states or undergo reconstruction of the configuration to give rise to non- defective structures or electronically benign defects75.

Surface- analysis techniquesStructural and chemical characterizationThe structures of surfaces are typically characterized at two distinct levels, atomic and morphological. As discussed above, the flexibility and vulnerability of the MHP lattice imparts rich chemistries and structural heterogeneities at the surfaces, profoundly affecting the macroscopic materials properties. Thus, there is high demand for scrupulous multi- perspective characteriza-tion techniques. Table 1 summarizes the surface sensitiv-ity and the pros and cons of the various characterization techniques employed in surface studies of MHPs.

Surface- structure- characterization techniques. Scanning electron microscopy (SEM) and topographic atomic force microscopy (AFM) are the most valuable and essential tools for characterizing the surface morphol-ogies of MHP thin films77–80. Both techniques provide information about the grain size and surface roughness, which are crucial parameters that govern the perfor-mance of thin- film devices81–84. These characterization methods offer complementary capabilities. SEM, com-bined with energy- dispersive X- ray spectroscopy, can measure the chemical composition of the surface structure, whereas there are numerous variants of AFM for the evaluation of other surface physical properties (discussed further below). Grazing- incidence X- ray scat-tering can be used to probe the lattice structure near sur-faces; surface sensitivity is enabled by adopting a small incident angle of the incoming X- ray. Incorporation of a 2D detector, as in grazing- incidence wide- angle X- ray scattering, enables the detection of the crystallographic orientation of the crystal lattice. By varying the incident angles, the penetration depths of the X- ray is modulated, enabling depth profiling of the lattice structure85,86. With an X- ray incident angle of 0.05° for near- surface detec-tion and 0.3° for full- depth detection, grazing- incidence wide- angle X- ray scattering signals determined the amount of PbI2 at the surface and throughout FAPbI3 thin films, revealing the formation mechanism of the FAPbI3 thin film86. Compared with the above techniques, higher spatial resolution, down to the atomic level, can be achieved using STM (Fig. 3a). In addition to some

Fig. 2 | surface chemistry and defects in metal halide perovskite nanoparticles. a | The various ligand- binding motifs according to the covalent- bond- classification method, illustrated here for a metal chalcogenide nanoparticle. b | Various ligands and surface- treatment methods have been adopted in the synthesis of perovskite nanoparticles. The ligands have been classified by their function: to improve the luminescent properties, enhance the electronic coupling or increase stability. The prototypical oleylamine–oleic acid ligand pairs used in the synthesis of metal halide perovskite quantum dots result in surface binding involving oleylammonium halide or oleylammonium oleate, as depicted in blue. c | Calculated defect- formation energies for interstitial, charged and neutral vacancies, and antisite substitutions on the surface of CsPbBr3 quantum dots (where C is core, SC is surface centre and SE is surface edge). VY denotes a Y vacancy, Yi denotes an interstitial Y- site and YZ denotes a Z- site substituted by Y, where Y and Z represent an ion of CsPbBr3. The right superscript denotes the charge of the defect and a left superscript denotes the number of the defects considered when calculating the formation energy. Backflip means that the antisite displacement goes back to its initial non- defective configuration. Calculations were performed using density functional theory with the PBE functional. Panel a adapted with permission from reF.65, ACS. Panel c adapted with permission from reF.75, ACS, further permissions relating to the material excerpted should be directed to the ACS.

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Table 1 | summary of surface- characterization techniques for metal halide perovskites

technique Property characterized resolution and detection depth

advantages limitations refs

Structural and chemical

SEM Topography Lateral: ~10 nm Can be combined with EDX for elemental analysis

Electron- beam damage; operates under UHV; information from two dimensions only

77,79,80

AFM Topography Lateral: ~10 nm

Depth: top surface

Multiple variants for evaluation of electronic properties; information from three dimensions

Lateral forces can distort the image

78

GIXS Lattice structure Depth: depends on incident angle, ~50–500 nm

GIWAXS can detect crystallographic orientation

Limited surface sensitivity 86

STM Surface atomic structure Lateral: atomic level

Depth: top surface

Highest spatial resolution Ultrasmooth surface required

87,88

HRTEM Lattice structure Lateral: atomic level High resolution; suitable for nanoparticles

Electron- beam damage; operates under UHV

89,94

XPS Surface chemistry Depth: ~10 nm Elemental sensitivity; quantitative Sample damage at UHV; limited chemical bonding information

15,96–98

ToF- SIMS Chemical composition Depth: depends on sputtering time, with a resolution of ~1 nm

Can produce an elemental depth profile

Destructive 104,105

Nano- FTIR Surface chemistry Lateral: ~10 nm

Depth: 10–100 nm

Molecular sensitivity Limited spatial resolution 112

SFG vibrational spectroscopy

Surface chemistry Depth: single molecular layer

Molecular orientation information; study of buried interfaces

Molecules in random orientations not detectable

113,114

Electronic and optical

ARPES Band dispersion Depth: 1–10 nm Electron distribution in momentum space for single- crystal studies

Limited to single crystals 116–118

UPS Electronic level below EF Depth: ~5 nm Precisely defines the position of the work function

Cannot reliably extract VBM 120,122–124

IPES Electronic level above EF Depth: ~1 nm Complementary to UPS Cannot reliably extract CBM 120

PL mapping PL spectra Lateral: ~1 μm Correlation between optical properties and surface microstructure when combined with SEM or AFM

Limited spatial resolution 145

cAFM Conductivity and topography

Same as AFM Distinct detection systems for topography and electronic properties

Contact mode; electronic properties of the bulk only

125–130

PFM Piezoresponse and topography

Same as AFM Simultaneously scans the topography and piezoresponse

Contact mode; piezoresponse of the bulk only

131,132

KPFM Surface EF and topography

Same as AFM CPD between tip and sample surface; photovoltage mapping; non- contact mode

Hard to obtain absolute value of EF without internal calibration; probe contamination

133–138

STS Local density of states Same as STM An extension of STM for electronic property measurement

Same as STM 35,88,139–141

TRS Surface carrier dynamics Depth: ~10 nm Surface recombination velocity can be extracted

Lack of detailed information on surface trap states

59,143

PAS Defect analysis Depth: depends on the energy of incident γ- ray

Can define the defect type Only sensitive to negatively charged defects

146,147

AFM, atomic force microscopy; ARPES, angle- resolved photoemission spectroscopy; cAFM, conductive AFM; CBM, conduction band minimum; CPD, contact potential difference; EDX, energy- dispersive X- ray spectroscopy; EF, Fermi level; GIWAXS, grazing- incidence wide- angle X- ray scattering; GIXS, grazing- incidence X- ray scattering; HRTEM, high- resolution transmission electron microscopy; IPES, inverse photoemission spectroscopy; KPFM, Kelvin probe force microscopy; nano- FTIR, synchrotron infrared nanospectroscopy; PAS, positron annihilation spectroscopy; PFM, piezoresponse force microscopy; PL, photoluminescence; SEM, scanning electron microscopy; SFG, sum- frequency generation; STM, scanning tunnelling microscopy; STS, scanning tunnelling spectroscopy; ToF- SIMS, time- of- flight secondary ion mass spectrometry; TRS, transient reflection spectroscopy; UHV, ultrahigh vacuum; UPS, ultraviolet photoelectron spectroscopy; VBM, valence band maximum; XPS, X- ray photoelectron spectroscopy.

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of the unique surface features of MAPbX3 discussed above, STM was used to pinpoint the exact locations of I− and Cl− in the mixed- anion MHPs MAPbBr3−xIx and MAPbBr3−yCly, revealing a disordered surface struc-ture with randomly distributed anion substitution and no phase segregation87 (Fig. 3b). The surface structural dynamics of MHPs can be probed in real time at the atomic level using photoexcited STM. For example, photoinduced reordering of the molecular dipole of MA+ was observed in MAPbBr3, and is speculated to be initiated by photoexcited separation of electron–hole pairs in spatially displaced orbitals88. However, ultra-smooth surfaces are a prerequisite for high- resolution STM measurements, requiring samples to be prepared through vacuum deposition or the cleavage of single crystals under ultrahigh vacuum, and thus limiting the application of this technique to a wider range of MHP systems, including solution- processed polycrystalline films. High- resolution transmission electron micros-copy (HRTEM) has also proved capable of atomic- scale detection of the outermost surface of MHP crystals. Using this technique, the thermally induced degrada-tion of MAPbI3 was investigated and shown to involve precipitation of trigonal PbI2 from an amorphized MAPbI3 layer89. Combined with density functional the-ory calculations, surface- initiated layer- by- layer degra-dation from tetragonal MAPbI3 to trigonal PbI2 was also observed using HRTEM90. However, the sensitivity of MHPs, especially their surfaces, to the electron- beam exposure can induce unexpected reactions and prod-ucts, leading to misleading conclusions42,91–93. MAPbI3, for example, undergoes continuous structural and com-positional changes with an increasing electron dose93. In this regard, a cryo- electron microscopy protocol94 was developed to mitigate, if not completely eliminate, the electron- beam damage to MHP samples. Additionally, HRTEM requires an electron- transparent specimen, usually with a thickness of within a few tens of nano-metres, to achieve high spatial resolution. Focused ion beam is a commonly used technique to prepare such thin samples, but damage to MHPs cannot be avoided. Direct deposition of a thin film on a transmission elec-tron microscopy grid prevents the perovskite specimen from being damaged, but the substrate dependence of the crystal growth can lead to variation in the properties of the thin films95. Although sample preparation poses an additional challenge to the application of HRTEM for investigating MHP bulk thin films, it is a go- to approach for the characterization of MHP nanoparticles, such as QDs12.

Surface- chemistry- characterization techniques. Numerous studies have used X- ray photoelectron spec-troscopy (XPS) to characterize the surface chemistry of MHP thin films, revealing the chemical composition and indicating molecular interactions based on the shift of the binding energy15,96–98. However, although XPS is advantageous for elemental analysis, it cannot precisely probe molecular bonding at an atomic level, particu-larly when the susceptibility of MHPs and the resulting radiolysis under ultrahigh vacuum are considered99. The ambiguous origin of the C 1s peak at ~284.8 eV,

presumably from MA+, FA+, solvent and/or adventi-tious carbon100–103, also contributes to the complexity of elucidating molecular structures with XPS. Time- of-flight secondary ion mass spectrometry can be used to identify the surface composition of MHPs through analysis of the ionized species sputtered from the sur-face of a specimen104,105. Additionally, this technique can provide a full- depth elemental- distribution profile when the sputtering proceeds to the bulk of the sam-ple, enabling evaluation of the variation in composition with depth106–110, although it still lacks information about molecular structure. Solid- state NMR spectroscopy has been used to unravel the atomic- level passivation mech-anism of the surfaces of MHP thin films upon treatment with various ammonium salts. However, although the molecular structures and the chemical environments of the organic ammoniums were determined, NMR techniques failed to establish whether these species are at the surface, owing to lack of surface sensitivity111. Synchrotron infrared nanospectroscopy (nano- FTIR) is a scanning- probe technique that takes advantage of the competence of Fourier transform infrared spectroscopy in exploring molecular structures and the superiority of AFM in acquiring spatial information. Nano- FTIR was recently applied to acquire spatial chemical informa-tion from MHP surfaces (Fig. 3c,d); spatial heterogeneity was detected in the vibrational activity of MHP films, which was ascribed to the nanoscale chemical diversity of isolated MHP grains112. Although nano- FTIR cannot provide a spatial resolution as high as that of STM, it succeeds in detecting molecular binding modes at the nanoscale without suffering from the limitations in sample- preparation methods, suggesting the importance of integrating characterization techniques for obtaining an in- depth understanding of the surface chemistry of MHPs.

The characterization techniques described above all fail to detect the molecular orientation of species at the surface; however, sum- frequency generation (SFG) vibrational spectroscopy can provide some insight. The SFG process cannot occur unless the inversion sym-metry of the medium is broken, which always takes place at surfaces or interfaces, making SFG vibrational spectroscopy a surface- sensitive and interface- sensitive technique (Fig. 3e). Subsequently, the molecular ori-entation at surfaces and interfaces can be determined by tuning the infrared beam to selectively probe the targeted molecule while controlling the polarization of each incident beam and the generated SFG beam113–115. For example, SFG vibrational spectroscopy was used to probe the molecular orientation of fullerene derivatives on the surface of MAPbI3 films113 (Fig. 3f). The polariza-tion combination was chosen so that the spectral inten-sity was proportional to the infrared dipole moment of vibrational modes with a component perpendicular to the substrate. The MAPbI3 films treated with different fullerene derivatives all showed peaks at ~1,463 cm−1, which corresponds to the vibrational mode of symme-try breaking in the fullerene derivatives with an infrared dipole moment pointing from the centre of C60 to the pendant group. The methylene symmetric stretching mode and methyl symmetric stretching mode of the

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pendant groups were also visible in the SFG spectra in the C–H stretch region. The sample treated with a fuller-ene derivative termed PCBB-3N-3I showed the strongest signals, indicating the assembly of PCBB-3N-3I on the surface with a preferred molecular orientation.

Electronic and optical characterizationUnderstanding the local electronic behaviour and optical properties of surfaces and the correlation with device performance is a crucial step towards the further development of MHP- based optoelectronic devices.

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Angle- resolved photoemission spectroscopy is typically used to determine the band dispersion of the surface of MHP single crystals in reciprocal space; single crystals are a good system for investigation, assuming the cleaved crystal is close to pristine with a minimally perturbed surface. The band mapping of the cleaved surfaces of MAPbI3 (reFs116–118) and MAPbBr3 (reFs26,117,119) single crystals is consistent with results from theoretical simu-lations. This agreement not only provides experimental support for the simulated electronic structure required to describe charge- carrier behaviour but also assists in

exploring and evaluating fundamental physical phenom-ena, such as Rashba splitting57. Ultraviolet photoelectron spectroscopy and inverse photoemission spectroscopy are complementary techniques for the direct and quan-titative measurement of the electronic- energy levels of MHP surfaces: the Fermi level and electron affinity can be extracted using ultraviolet photoelectron spectroscopy, whereas inverse photoemission spectroscopy probes the electronic states above the Fermi level of the system120 (Fig. 3g,h). Although these techniques have been widely adopted for the measurement of band edges121–124, the soft energy onsets of MHP surfaces, originating from the structural disorder and low density of states at band edges, pose a challenge for the accurate determination of the energy levels, causing variations in the reported values. All these approaches for characterization of the electronic structure provide ensemble surface properties of MHPs without giving local spatial information.

Various scanning- probe- microscopy techniques, var-iants of topographic AFM or STM for probing surface atomic structure and morphology, have been widely used to explore local energetics on MHP surfaces at the nano-scale or atomic scale. With conductive AFM (cAFM), the conductivity can be mapped by measuring the current between a conductive tip and substrate and then used to assess local electrical properties, such as charge distri-bution or transport125–128. The correlation between local photocurrent characteristics and surface topography can be probed with an additional light stimulus shining onto the sample surface. For example, intragrain heterogene-ity, indicative of facet- dependent defect concentrations, was identified in a polycrystalline MHP thin film using cAFM under luminescence129. Combined with confocal photoluminescence (PL) and laser- beam- induced cur-rent, illuminated- cAFM measurements unravelled an anticorrelation between the optical and photoelectric response of a MHP thin film130. As a variant of cAFM, piezoresponse force microscopy can aid in determining the morphology and, simultaneously, the piezoresponse of MHP surfaces through application of an alternating voltage to the probe. A good example of its success in providing insight into the unique physical properties of MHP surfaces is the experimental evidence that shows the existence of ferroelastic domains in MAPbI3 (reFs131,132). By contrast, Kelvin probe force microscopy (KPFM) probes the contact potential difference between a tip and a sample surface, from which the surface work function can be extracted if calibrated carefully124,133–135. Illuminated- KPFM also offers the opportunity to map the open- circuit voltage when the back contact layer of  the MHP is grounded with respect to the tip136. Advances in KPFM have enabled a fast scanning mode on the timescale of seconds, allowing for real- time inves-tigation with a nanoscale spatial resolution137. However, when conducted under vacuum, supraband- gap light irradiation induced the loss of halide species, which adsorb on the Kelvin probe tip138. If undetected, this can lead to misinterpretation of the MHP surface potential. Scanning tunnelling spectroscopy, an extension of STM, provides the local density of states with a high spatial and energetic resolution. By ramping the bias voltage, both the occupied and unoccupied states can be obtained,

Fig. 3 | surface-sensitive techniques for characterizing metal halide perovskites. a | Scanning tunnelling microscopy (STM). When a conductive tip is brought close to a sample surface, an applied bias between them results in the tunnelling current, which is a function of the tip position, applied voltage and the local density of states of the sample. b | STM images of MAPbBr3, MAPbBr3−xIx and MAPbBr3−yCly (where MA+ is methyl ammonium) surfaces (top) and the corresponding large- area STM images (10 nm × 10 nm; bottom). The bright and dark protrusions observed in the mixed- halide perovskite STM images were assigned, respectively, to I− and Cl−, which substitute Br− at the surface of the perovskite films. The insets show fast- Fourier- transform results obtained from the topographic STM images; compared with MAPbBr3, the mixed- halide perovskites do not show additional peaks at low k- values (where k is the wavevector), suggesting the long- range disorder of the additional I− and Cl− at the surface. c | Atomic force microscopy (AFM) and AFM- related surface- characterization techniques. When the tip gets very close to a sample surface, the cantilever is deflected by the tip–surface interaction. Different physical quantities can be measured using specialized probes under various external stimuli. For example, in synchrotron infrared nanospectroscopy (nano- FTIR), the tip is illuminated by an external light source (IR synchrotron) and the tip- scattered light is detected as a function of tip position. d | AFM topography map of a CsFAMA- based perovskite (where FA+ is formamidinium; left), the corresponding IR broadband image (middle) and the nano- FTIR point spectra from the numbered regions (right), revealing vibrational heterogeneity. e | Sum- frequency generation (SFG) vibrational spectroscopy. Visible and IR beams (frequency tunable or broadband) are typically used as input beams (where ωVIS and ωIR are the frequencies of the visible and IR input beams, respectively). An SFG signal (with a frequency ωSFG) is generated from a surface or interface at which inversion symmetry is broken. f | SFG spectra (top) of MAPbI3 before and after treatment with various fullerene derivatives (phenyl- C61- butyric acid methyl ester (PCBM) and two other derivatives called PCBB-3N-3I and PCBB-3N) at the regions corresponding to the vibrational modes of C60 (left) and the C–H stretching region of the pendant groups (right). Schematics (bottom) of the molecular orientation of PCBB-3N-3I (left) and PCBB-3N (right) on MAPbI3. PCBB-3N-3I assembles with a molecular orientation that is more ordered than that of PCBB-3N, and thus shows a stronger signal in the SFG spectra. g | Photoelectron spectroscopy. The energies of the emitted photoelectrons, which can escape only from a depth of a few nanometres, are characteristic of their original electronic states. h | Representative ultraviolet photoelectron spectroscopy (UPS) and inverse photoemission spectroscopy (IPES) spectra of 18 metal halide perovskites. For better comparability, the curves are offset vertically and the high- energy cut- offs are aligned at an excitation energy of 21.22 eV, marked by line 1. Lines 2, 3 and 4 indicate characteristic features in the density of states, corresponding to the position of states related to Cs+, MA+ and FA+, respectively. The extracted positions of the valance band maximum (VBM) and conduction band minimum (CBM) are indicated by black vertical markers, and the Fermi level (EF) positions are marked by triangles. i | Transient reflection spectroscopy. In transient reflection spectroscopy, the sample is excited with a short laser pump pulse, followed by a time- delayed white- light continuum probe pulse that spans the bandgap of the semiconductor surface of interest. The white- light continuum reflects off the sample surface and the pump- induced change in reflectance (ΔR/R) is measured. j | The normalized surface carrier kinetics of polycrystalline perovskite films photoexcited at different photon energies. cAFM, conductive AFM; KPFM, Kelvin probe force microscopy; PFM, piezoresponse force microscopy. Panel b adapted with permission from reF.87, ACS. Panel d adapted with permission of AAAS from reF.112, © The Authors, some rights reserved; exclusive licensee American Association for the Advancement of Science. Distributed under a Creative Commons Attribution NonCommercial License 4.0 (CC BY- NC) http://creativecommons.org/licenses/ by- nc/4.0/. Panel f adapted from reF.113, CC BY 4.0. Panel h adapted from reF.120, CC BY 4.0. Panel j adapted from reF.59, Springer Nature Limited.

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serving as a powerful technique for the determination of surface band structure35,88,139–142.

By investigating the photon- generated carrier dyna-mics, the surface defects of MHPs can be evaluated. Transient reflection spectroscopy (Fig. 3i) has proved to be a powerful tool for measuring surface carrier dynamics, from which the surface carrier- diffusion coefficient and surface recombination velocity (SRV) can be derived59,143,144 (Fig. 3j). Comparing the SRV of a MHP single crystal and polycrystalline film revealed unintended surface passivation of the polycrystalline thin film, indicated by the nearly an order of magni-tude smaller SRV than in its single- crystal counterpart. Despite the relatively low SRV, surface recombination, rather than intergrain recombination, was still sug-gested to be dominant in regulating the carrier dynamics in the polycrystalline film, emphasizing the impor-tance of intentional surface- passivation strategies in MHP- based optoelectronic devices59. When combined with SEM or AFM, PL mapping of MHP surfaces can be achieved using confocal PL microscopy to correlate the morphological characteristics with the local opti-cal properties145. Positron annihilation spectroscopy is also a surface- sensitive and depth- sensitive technique for studying the defects in MHPs by varying the energy of the incident positrons for control of the penetration depth. As the positrons are injected into the sample, the positron–electron annihilation process, which proceeds by the emission of γ- rays, yields information on the local electron- density momentum, reflecting the defect type, distribution and concentration within the sample146,147. However, this technique is sensitive only to open volume defects, and positrons fail to ‘observe’ positively charged defects, deeming it limited for comprehensive analysis of surface defects.

Surface manipulationSurface role in bulk materials and devicesThe most pressing question remaining in the context of improving device performance is how to achieve the optimal materials properties by manipulating the sur-face of MHPs. Owing to the large number of defects at the surface that serve as carrier- recombination sites, the most straightforward and widely adopted strategy is to deactivate these surface traps. The ionic nature of MHPs enables defect- healing approaches based on Lewis acid–base chemistry. Undercoordinated Pb2+ or Pb clusters on the surface act as Lewis acids and can accept non- bonding electrons from a Lewis base to deactivate their carrier- trapping capability. Molecules with atoms bearing lone pairs of electrons, such as oxygen15,97,148–151, nitrogen7,121,148,152–154, sulfur98,121,155 and phosphorus156,157, are thus exploited as Lewis bases to donate electrons to these surface defects. The extent to which such passiva-tion molecules minimize surface non- radiative recom-bination losses is dependent on the interaction with the defect sites (Fig. 4a). For a series of passivating molecules with amino groups, the introduction of oxygen atoms weakened the hydrogen bonding between the amino groups and perovskite organic cations, enhancing coor-dination with the Pb2+ defect sites. Stronger coordina-tion with the Pb2+ sites increases the external quantum

efficiency of LED devices. In particular, the formation of strong Pb–F (reF.105), Pb–Cl (reF.158) and Pb–O (reFs158,159) bonds or the formation of a robust lead oxysalt species on the surface160 has proved to be beneficial to both the efficiency and the stability of MHP devices. For example, the in situ formation of a lead sulfate surface layer on a MHP thin film (Fig. 4b, top) resulted in a reduction in the trap density of states within photovoltaic devices (Fig. 4b, bottom), indicating enhanced PCE and stability161.

Undercoordinated halide trap sites are Lewis bases and can be silenced by binding with various Lewis acids, among which fullerene and its derivatives (such as phenyl- C61- butyric acid methyl ester (PCBM)) have been most intensively investigated160,162,163. Although the passivation mechanism involving fullerene species is not yet fully understood, such as the binding site with MHPs, early studies point to the formation of a PCBM–halide radical, as well as wavefunction hybridization between PCBM and the perovskite surface164. In addition to deal-ing with Pb- based and halide- based defects, the impor-tance of manipulating the surface organic cations is evidenced by the inhibition of FA+ vacancies through the formation of hydrogen bonds between FA+ and F− (reF.105) (Fig. 4c) and the diboron- assisted decomposition of unre-acted organic salts on the surface to decrease the defect density165. As the MHP lattice is ionic, its surface defects are charged, and, thus, electrostatic interactions also have a role in surface- defect passivation. Ammonium cations have been explored extensively for interaction with surface defects, mainly through electrostatic inter-actions. Under high- temperature annealing conditions, the organic cations are lost from the surface of MHP thin films at appreciable rates166. Added ammonium cations, therefore, interact with the remaining Pb2+ and X− on the surface. Surface passivation with organic ammo-niums typically results in enhanced PL (Fig. 4d, top), which suggests a decrease in non- radiative recombina-tion sites and, thus, improved photovoltaic PCEs (Fig. 4d, bottom). Although various ammoniums, including MA+ (reFs9,167–169), ethylammonium111, butylammonium170–172, tetrabutylammonium173, octy lammo nium97,104,174, phenyl trimethylammonium175, pheny l ethylammon-ium19,123,135,176–182 and diammoniums133,183,184, have been reported to be effective for MHP surface treatment and increasing the efficiencies of devices, the precise passivation mechanism remains ambiguous. This ambi-guity is likely due to the formation of 2D perovskites and the modification of the band alignment at the surface, making it difficult to attribute the passivation effects solely to the electrostatic interaction. Additionally, bear-ing in mind the complexity of charged surface defects, zwitterions (which have both positive and negative charges) have been shown to be passivation agents for surface trap healing. Capable of interacting with both negatively and positively charged defects, zwitterions such as l- α- phosphatidylcholine have demonstrated the significance of considering charge neutrality when manipulating the surface defects of MHPs185,186.

Trap healing on MHP surfaces is usually accompa-nied with a change in the Fermi energy or band- edge position, as a result of the absence of Fermi- level pin-ning by trap states or electron- orbital redistribution at

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the surface. The shift in the energy levels provides evi-dence of surface- defect passivation, but also induces favourable band alignment, giving rise to enhanced carrier transport. Introducing a surface dipole can also adjust the band structure and improve carrier dynamics. A surface dipole can be introduced through a passivating molecule with a strong intrinsic electric dipole moment located on the surface113,124,187–189 or the redistribution of electron density owing to the interaction of a molecule with the MHP surface167,190–192. Both approaches build an additional electric field into the interface that mod-ulates the energy levels of a MHP surface to adjust the

band alignment with the transporting layer and, thus, facilitate carrier injection or extraction. As an example, band bending at a perovskite surface can be tuned by switching the direction of the dipole moment of mole-cules at the surface (Fig. 5a). The band structure at the surface of MHPs can also be modulated by forming a heterostructure through the incorporation of a second semiconductor or insulating capping layer. It is possible to decrease surface carrier recombination and enhance carrier extraction by constructing graded energy levels towards selective contact134,171,176,193–195 (Fig. 5b, left) or by capping a wide- bandgap layer (Fig. 5b, right), enabling

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tunnelling of the minority carrier and blocking the majority carrier172,191,196–199.

In addition to regulating the surface electronic structure through trap healing and band tuning, the surface energy provides an additional handle through which the  surface properties can be manipulated. Most notably, the surface tension can be decreased to yield a hydrophobic surface that is resistant to external moisture. Various hydrophobic polymers and organic molecules have been used to develop MHP films with a low surface tension, thus prohibiting water penetra-tion and increasing device stability173,180,182,200–205 (Fig. 6a). Surface- energy- driven grain growth or recrystallization is a unique aspect of MHP surface manipulation, leading to reconstruction of the entire bulk structure85,104,167,169. By modifying the surface of a polycrystalline thin film to lower the surface energy of a specific crystallographic facet, the crystalline grains subsequently preferentially grow along that orientation to minimize the free energy of the system. Therefore, the grain size of MHPs can be rationally controlled by varying the surface- energy anisotropy with different surface- treatment agents. For example, a series of organic ammoniums induced an increase in surface- energy anisotropy in the order butylammonium < octylammonium < oleylammonium, resulting in a corresponding increase in the size of the secondary grains with a preferred orientation104 (Fig. 6b). The soft lattice of MHPs is a particularly crucial factor in facilitating this surface- induced structural evolu-tion over the entire bulk thin film; by contrast, such surface effects are dominant only at the nanoscale in conventional semiconductors104.

Surface- modification strategies such as trap healing, band tuning or surface- energy manipulation ultimately improve the interfacial properties in MHP- based pho-tovoltaic or LED devices. Decreased defect density, improved band alignment and preferred crystallographic orientation improve the interfacial- energy landscape, facilitating carrier transport and, thus, increasing the efficiency of as- fabricated devices. As device degrada-tion always initiates at an interface, it is beneficial to the long- term stability to eliminate surface structural disor-der and lower the surface tension to prevent moisture penetration.

Surface role in nanomaterials and devicesThe unusually rapid success in producing MHP QDs with near- unity PLQYs owing to their unique defect- tolerant photophysics seems to suggest that it is less important to eliminate surface defects than in conventional semi-conductor nanoparticles. Despite the incredible success, a key question is: what prevents the PLQY from being unity? Hence, efforts have been devoted to determining the surface states of MHP nanoparticles and explor-ing new surface chemistries beyond the prototypical oleylamine–oleic acid ligand pair to further decrease the number of surface traps and improve optical or optoelectronic properties. In Cs+- based MHP nano-particles, a halide self- passivation effect was proposed upon obtaining a high quantum yield with a halide- rich composition206–208. This finding is consistent with theo-retical results (discussed above) that demonstrated that AX- terminated MHP QDs are more defect tolerant than PbX2- terminated MHP QDs. Similar to their bulk coun-terparts, various Lewis bases, such as alkylphosphonic acid, trioctylphosphine, oxygen and thiocyanate209–212, decrease or eliminate surface trap states in MHP QDs and, thus, improve their luminescent properties (Fig. 2b, luminescent properties). Using hard–soft acid–base theory to guide their design, soft anionic Lewis bases were used to target undercoordinated Pb2+ and produce essentially trap- free CsPbX3 QDs213. Surface- defect passivation is especially important when dealing with the ligand- stripping process. The presence of residual long- chain organic molecules and ligands on the sur-face of QDs increases the inter- dot spacing and pre-vents electronic coupling in QD solids; the immediate solution to this problem is to strip the ligands or replace them with shorter ones14,214. The dynamic ligand bind-ing in MHP QDs causes the ligands to quickly desorb, creating surface traps. Therefore, handling the surface ligands without sacrificing the optoelectronic proper-ties or affecting the structural integrity poses a chal-lenge to developing highly conductive MHP QD films. Several attempts have been made in this regard, includ-ing post- treatment with ammonium halides or metal halides215–217, solvent engineering70,218,219, crosslinking220 and the adoption of a substitutional ligand. Among these strategies, developing ligands that maintain the struc-tural integrity without hindering carrier transport has attracted considerable attention. Typically, ligands with shorter alkyl chains206,221, a conjugated structure222–224 and zwitterionic features225 have been exploited (Fig. 2b, electronic coupling).

Fig. 4 | surface-defect-healing strategies. a | Molecular structures and calculated electron distributions of various amine- based passivation agents (top). The average peak external quantum efficiency (EQE) of light- emitting diodes based on metal halide perovskites treated with these agents is dependent on the adsorption energy difference, ΔEad = Ead,V − Ead,P, where Ead,P and Ead,V are the absorption energies of the passivation agent on a perfect surface and a surface with defects, respectively (bottom). Each value is an average of 60 devices, and the error bars correspond to the standard deviation. The calculated electron distribution of HMDA and EDEA (top) show how the electrons at the nitrogen atoms are inductively polarized towards the oxygen atoms, which decreases the hydrogen bonding with organic cations but increases the interaction with the surface defects. b | Perovskite surfaces can be passivated through in situ formation of a lead sulfate top layer (top, where MA+ is methylammonium and FA+ is formamidinium). Comparison of the trap density of states (tDOS) obtained through thermal admittance spectroscopy as a function of trap energy E(ω) for devices based on perovskites with or without lead sulfate top layers (bottom). The lead sulfate top layer decreases the tDOS over nearly all trap- energy depths. The vertical dashed lines define trap bands with different trap- energy depths: shallow trap states (band 1) and deep trap states (band 2 and band 3). c | Optimized structures of a Na–I unit and a Na–F unit adjacent to an FAPbI3 surface (top). The FA+ cations surrounding an F− site reorient towards F− to maximize their interaction. Comparison of the calculated formation energy of a surface FA+ vacancy in clean FAPbI3 (FA) and in surfaces with the NaI (FA–I) or NaF (FA–F) incorporated (bottom). Incorporation of F− increases the formation energy by a sizeable 0.55 eV, suggesting that F− helps to prevent the formation of FA+ vacancies. d | Steady photoluminescence (PL) of FA1−xMAxPbI3 perovskite films with PEAI (where PEA+ is phenylethylammonium) treatment under different conditions (top) and typical current–voltage curves of devices based on films with and without PEAI treatment under 1- sun conditions (bottom). DDDA, 4,9-dioxa-1,12- dodecanediamine; EDEA, 2,2′-(ethylenedioxy)diethylamine; HMDA, hexa-methylenediamine; ODEA, 2,2-[oxybis(ethylenoxy)]diethylamine; RT, room temperature; TTDDA, 4,7 ,10- trioxa-1,13- tridecanediamine. Panel a adapted from reF.7, Springer Nature Limited. Panel b adapted with permission from reF.161, AAAS. Panel c adapted from reF.105, Springer Nature Limited. Panel d adapted from reF.19, Springer Nature Limited.

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The effects of the surface on MHP stability become even more predominant at the nanoscale. Although cubic- phase CsPbI3 and FAPbI3 bulk materials suffer from thermodynamic instability at room tempera-ture, folding them into QDs stabilizes the photoactive phase218,226,227. This stabilization is attributed to the enlarged contribution from the surface energy in QDs, as well as the strain redistribution through the lat-tice owing to size effects218. Like their bulk analogues, MHP QDs exhibit fast degradation upon exposure to external stimuli, such as moisture or other polar envi-ronments. The large surface- to- volume ratio in QDs further facilitates this process. Various strategies have been rapidly developed to modify the surface of MHP QDs with a hydrophobic protective layer, such as silica or

polymers227–236 (Fig. 2b, stability). Embedding MHP QDs into a robust crystalline inorganic structure was recently proposed to be an effective approach to increasing the stability without notably sacrificing optoelectronic prop-erties. Examples of the inorganic structures include the 2D perovskite OA2PbBr4 (where OA+ is octylammo-nium)237, 0D perovskite Cs4PbBr6 (reF.238) and PbSO4 (reF.239); each has compositional or structural features in common with the original MHP QDs, with the aim of establishing surface compatibility. The limited stability of MHP QDs has also been attributed to facile proton exchange between the oleate and amine surfactants, as well as to amine- accelerated dissolution of lead oleate from the surface. These suggestions have inspired the exploration of amine- free surface chemistry. MHP QDs

–5.7 eV

a

b

MHPTiO2

HS–Ph–CN

MHP HTL MHP HTL MHP

MHP

Insulatinglayer

DOS

C60

HTL

HTL MHPTiO2

HS–Ph–SCH3

HTL

CBM

VBM

–5.7 eV

–5.2 eV

–5.6 eV

–5.8 eV

Without With NCs With NWs

CBM (–3.9 eV)

VBM (–5.4 eV)

LUMO(–4.5 eV)

HOMO (–6.2 eV)

+

–

+

+

–

+

+

+

+

+

+

––

––

––

–

–

–

––– –

–––

–

(Cs,Br)-richperovskite

–5.2 eV

Electrondensity

Holedensity

Hole

Electron SCH3

SH

CN

SH

CBM

VBM

Fig. 5 | surface band-tuning strategies. a | The energy levels at the interface between a metal halide perovskite (MHP) and a hole- transporting layer (HTL) under open- circuit conditions are altered by modification of the interface with benzenethiol molecules, which have a notable dipole moment. With a CN- substituted benzenethiol (HS–Ph–CN), the molecular dipole points away from the MHP, leading to an upwards shift in the surface- energy level of the MHP (left). By contrast, a downwards shift is observed when the interface is modified with a SCH3- substituted benzenethiol (HS–Ph–SCH3) and the molecular dipole points towards the MHP (right). b | Schematic representation of hole extraction in MHP solar cells without nanocrystals, with CsPbBr3 nanocrystals (NCs) and with CsPbBr3 nanowires (NWs) (left). Without NCs, there is a notable energy- level mismatch between the HTL and the valence band maximum (VBM) of the MHP. Incorporation of NCs increases the VBM by 0.09 eV, improving the band alignment with the HTL and, thus, enhancing carrier extraction and the device efficiency. Compared with NCs, the incorporation of NWs makes the VBM increase more uniformly without forming mixed electronic states on the surface, further enhancing device efficiency. The energy diagram (right) illustrates the suppression of surface- charge recombination by the introduction of an insulating layer. Matching the energy of the conduction band minimum (CBM) to the unoccupied states in the electron- transporting layer (C60), enables the electrons in the MHP to tunnel into the latter. By contrast, the holes in the MHP are blocked, leading to spatial separation of the electrons and holes and, thus, a decrease in charge recombination. DOS, density of states; HOMO, highest occupied molecular orbital; LUMO, lowest unoccupied molecular orbital. Panel a adapted with permission from reF.124, RSC. Panel b adapted with permission from (left) reF.194, Wiley and (right) reF.191, Wiley.

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stabilized solely by either benzenesulfonic acid240 or oleic acid derived from tetraoctylammonium- halide precursors241 have been synthesized. The enhanced stability against polar solvents used during purifica-tion deems this approach a promising pathway towards high- performance QD- based optoelectronic devices.

Future perspectivesDespite extensive developments in the passivation of surface defects, a comprehensive understanding of the origins of surface states in MHPs remains to be achieved. This knowledge gap emphasizes the significance of meticulously identifying surface defects for the rational design of surface- passivation approaches without trial and error. As most of the current surface- defect analy-ses are based on theoretical modelling, experimental studies are direly needed to pinpoint the types, energy

levels, formation conditions and locations of the defects on perovskite surfaces. Considering the complexity of the surface- defect configuration, it is challenging to gain a complete picture with a single technique. Introducing various molecules known to target specific defects could help to unveil the type of surface defects present and their atomistic origins. The design of molecular struc-tures for targeted passivation combined with a set of surface- sensitive techniques for evaluating the passi-vation efficacy would provide further insight into the surface states.

Long- term stability of MHPs remains a major chal-lenge, requiring new techniques that enable in situ and low- damage or damage- free surface characterization. In situ analysis would enable the real- time monitoring of the evolution of properties within operational MHP devices, aiding in determining the source of stability

0 200 400 600 800 1,000

3D/2D PSC3D PSC

1.0

0.5PCE

Time (hours)

a b0 s

1 s

2 s

3 s

0 s

1 s

2 s

3 s

96.0°

46.8°

44.2°

40.0°

38.4°

96.0°

95.8°

95.8°

Large grains and preferred orientationSmall grains and random orientation

γmin

SISG

γ

BA (100)γ

100 = 3.23

OA (100)γ

100 = 2.14

OLA (100)γ

100 = 1.97

BA (111)γ

111 = 4.09

OA (111)γ

111 = 4.06

OLA (111)γ

111 = 4.01

Cs Pb I Br C N H

400 nm

R= NH3

Ref BA OA OLA

Fig. 6 | Manipulation of surface tension. a | Images of water droplets on the surface of 3D/2D (top, left) and 3D (top, right) perovskite films at different water- loading times. The 2D perovskite contains the pentafluorophenylethylammonium cation, which is hydrophobic and imparts increased moisture resistance. The ambient- atmosphere (40% relative humidity) ageing results of a 3D and a 3D/2D perovskite solar cell (PSC) at 1- sun irradiation show that the 3D/2D device maintains better long- term operational stability (bottom). b | Optimized (100) and (111) slab models of a perovskite with various ammonium surface modifications (top). The difference in the surface energy of the (100) and (111) planes (γ100 and γ111, respectively, in units of eV nm−2) is butylammonium (BA+) < octylammonium (OA+) < oleylammonium (OLA+), indicating a corresponding increase in surface- energy anisotropy. The top- view scanning electron microscopy images (middle) of perovskite films with various treatments and a reference system (Ref) without treatment show that surface treatment increases the size of the grains, with a correlation between the average grain size and the surface- energy anisotropy. The increase in grain size is attributed to surface- induced secondary grain growth (SISG; bottom). The driving force for SISG is the surface- energy anisotropy, defined as γ γ γ∆ = −

min, where γ is the average surface energy and γ

min is the

minimum surface energy. Upon incorporation of ammonium cations at the surface, a specific crystallographic facet shows γ

min, and the grains prefer to grow along that orientation. PCE, power- conversion efficiency. Panel a adapted with

permission of AAAS from reF.182, © The Authors, some rights reserved; exclusive licensee American Association for the Advancement of Science. Distributed under a Creative Commons Attribution NonCommercial License 4.0 (CC BY- NC). Panel b adapted with permission from reF.104, ACS.

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issues during operation. The fragile lattice of perovskites is highly sensitive to beam sources, adding to the dif-ficulties in precisely determining the material proper-ties of the surface region. For example, HRTEM fails to capture atomic- resolution images of organic–inorganic hybrid MHPs, owing to electron- beam damage. A signif-icant variation in the elemental composition measured using XPS has also been observed. It is hard to unambig-uously answer whether the discrepancies are a result of different processing histories or the X- ray damage to the samples. Recently, cryo- HRTEM successfully acquired high- resolution MHP structural information with min-imal damage from the electrons94. Sample transfer and measurements under cryo conditions minimizes struc-tural transformations, making it a promising approach towards clear- cut determination of atomic- level con-struction. However, cryo conditions cannot entirely eliminate electron- beam damage, especially for some highly volatile organic components in MHPs, and, therefore, an electron dose test is also needed.

Although various surface- passivation agents improve MHP materials properties or device performance, screening a tremendously large molecule pool poses a big challenge. For example, molecules containing a carbonyl group, a well- known Lewis base, can act as passivation agents for perovskites. However, although the rich chem-istry in organic molecules provides structural tunability, it increases the complexity of finding the most effective passivating molecules. There are myriad molecules con-taining carbonyl groups; however, the molecular struc-ture that will be most effective is often difficult to judge. A recent breakthrough in circumventing this problem involved the systematic investigation of the chemical environment around the effective functional group and its effects on surface passivation15. This breakthrough

should motivate further efforts towards rational molecular design for surface passivation.

In addition to molecular surface- passivation approach es, a thin surface layer of a robust inorganic material that not only silences the surface defects but also readily binds to the soft perovskite surface is highly desirable for long- term stability. The success of PbO or PbSO4 surface layers on pristine MHPs158,159,161, mim-icking the SiO2 surface layer used in Si photovoltaics, implies that this strategy should become one of the major research directions for surface passivation in MHP optoelectronic devices.

The structural or compositional similarity is a prereq-uisite to epitaxial construction of a surface- passivation layer. Therefore, in addition to Pb derivatives, various organic ammonium halides are also intensively studied for surface modification in MHPs. Although the under-lying passivation mechanism remains ambiguous, the capability of achieving superior device performance is encouraging. Additionally, the organic tail of the ammo-nium cations can be functionalized to modify the surface properties, which offers enormous scope to tune the band structure and, thus, carrier dynamics at MHP surfaces.

Last, but not least, the large surface- to- volume ratio of QDs may emphasize the surface states; thus, QDs could be exploited as model systems for the study of perovskite surfaces. Although there are significant dif-ferences between the surface of the bulk and that of the QDs, they share common characteristics, as evidenced by a few similar surface- passivation methods reported for both QDs and bulk MHPs156,211,217,242. Successful surface- passivation strategies in perovskite QDs can be applied to perovskite bulk films or vice versa.

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AcknowledgementsThis article is based on work supported by the US Department of Energy’s Office of Energy Efficiency and Renewable Energy (EERE) under the Solar Energy Technologies Office award num-ber DE- EE0008751. The authors thank A. Z. Stieg and S. Gilbert Corder for helpful discussions regarding AFM and nano- FTIR techniques, and S. Nuryyeva, M. Mujahid and S. Tan for help with language editing prior to submission.

Author contributionsJ.X. and R.W. researched most of the data and prepared the first draft. Y.Y. revised the manuscript before submission and supervised the project.

Competing interestsThe authors declare no competing interests.

Publisher’s noteSpringer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. © Springer Nature Limited 2020

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