The role of grain boundaries in perovskite solar cellsyylab.seas.ucla.edu/papers/MTE-review.pdfThe role of grain boundaries in perovskite solar cells Jin-Wook Lee 1, Sang-Hoon Bae

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Materials Today Energy xxx (2017) 1e12

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Materials Today Energy

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The role of grain boundaries in perovskite solar cells

Jin-Wook Lee 1, Sang-Hoon Bae 1, Nicholas De Marco 1, Yao-Tsung Hsieh 1, Zhenghong Dai,Yang Yang*

Department of Materials Science and Engineering and California NanoSystems Institute, University of California, Los Angeles, CA 90095, United States

a r t i c l e i n f o

Article history:Received 15 May 2017Accepted 26 July 2017Available online xxx

Keywords:PerovskiteGrain boundaryDefectStabilityPassivation

* Corresponding author.E-mail address: [email protected] (Y. Yang).

1 These authors contributed equally to this work.

http://dx.doi.org/10.1016/j.mtener.2017.07.0142468-6069/© 2017 Elsevier Ltd. All rights reserved.

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a b s t r a c t

With a certified power conversion efficiency (PCE) exceeding 22%, the hunt is now on clues to furtherimprove the PCE and stability of perovskite (PVSK) solar cells toward commercialization. PolycrystallinePVSK films are grown by low temperature solution processes that results in the formation of sub-micronscale grains and consequent grain boundaries. The large discrepancy observed between optoelectronicproperties of the polycrystalline films and single crystals of PVSK implies that grain boundaries maylargely influence the optoelectronic properties of the PVSK thin film. In this article, important studies onthe roles of the grain boundaries on optoelectronic properties and stability of the PVSK are reviewed.

© 2017 Elsevier Ltd. All rights reserved.

1. Introduction

The consistent efforts to find a silicon alternative for solar cellsbrought about a game changer in 2009 with the perovskite (PVSK)light absorber [1,2]. With an important milestone in 2012 byintroducing the solid-state version of PVSK solar cells, greatattention has been focused on PVSK solar cells [3,4] yielding acertified power conversion efficiency (PCE) of 22.1% by 2016 [5].Benefitting from low temperature and scalable solution processes,the PVSK solar cell is now one of the most promising candidates toreplace the silicon solar cells [6e8].

The unprecedented rapid progress in PCE of PVSK solar cells hasbeen attributed to its superb optoelectronic properties. The PVSKmaterials have a high absorption coefficient, long-range balancedelectron and hole diffusion lengths and high defect tolerance,enabling efficient charge collection with comparable open-circuitvoltage (VOC) deficit with single crystal GaAs solar cells [9e13].With the consideration of photon recycling effect, PCE of 31% wassuggested to be achievable for the PVSK solar cells [14].

Accompanied with remarkable progresses in film formation andcompositional engineering [15e21], researches have started look-ing into the microscopic properties of the films to find the room forfurther improvement [22e24]. As the PVSK thin films are grown by

et al., The role of grain bound

low temperature solution processes, the formation of poly-crystalline films with grain boundaries (GBs) is inevitable. As aresult, the structural disorders at GBs potentially can form defectstates, which possibly degrade the optoelectronic quality of thePVSK films. Indeed, the charge carrier diffusion lengths measuredfrom polycrystalline thin film and single crystal PVSK show largediscrepancy [10,12,25,26], which implying clues for furtherenhancement of PCE possibly lays on the microstructure.

Recent studies on local proprieties on PVSK have revealed thatthe GBs and defects associated them have important roles in bothphotovoltaic performance and stability of the PVSK solar cells[23,27e30]. In this article, we introduce the studies on micro-structure using a variety of analytical tools including computation,photo-physical characterization and atomic force microscope(AFM). In particular, from crystallization and formation of GBs inPVSK films, to their influences on optoelectronic and photovoltaicproperties will be addressed. Finally, important progresses topassivate the GBs will be reviewed.

2. The roles of GBs in PVSK solar cells

2.1. Phase formation and crystal growth of the PVSK

The PVSK materials used for photovoltaic applications employthe ABX3 crystal structure, inwhich A cations in cubo-octahedral sitemake bonding with BX6 octahedra. According to the tolerance factorwhich provides the geometrical guideline of formability for thePVSK structure, various components are available to sustain the

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J.-W. Lee et al. / Materials Today Energy xxx (2017) 1e122

PVSK structure [31]. The relative sizes of the A cation and the BX6octahedra determine the dimensionality of the PVSK materials. ThePVSK material can remain three-dimensional (3D) with relativelysmall ‘A’ cations while incorporation of larger cations result in two-dimensional (2D) or even one-dimensional (1D) chain materialswith formulas A2BX4 and A3BX5 [32]. For photovoltaics, the mostrepresentative components for each site are methylammonium ion(CH3NH3

þ, MAþ) and formamidinium ion (CH5N2þ, FAþ) for ‘A’ site,

Pb2þ and Sn2þ for ‘B’ site I�, Br�, and Cl� for ‘X’ site. Among severalcombinations of the components, the most common halide PVSK isMAPbI3. In most of cases, the reaction for phase formation followssimple chemical equation (e.g. PbI2 þ MAI / MAPbI3) withremarkably quick reaction rate [33]. In case of MAPbI3, the phasetransformation is relatively simple, but in the mixed compositions,such as mixed cation, mixed-group IV metal, and mixed-halidePVSKs, it becomes more complicated. For the mixed system, extraeffects coming from different elements such as size effect, chemicaleffect should be taken into account. For example, addition of chlo-ride (Cl�) to theMAPbI3 PVSK enables tuning of electrical and opticalproperties, so it drew an attention at the beginning of halide PVSKresearch [34]. However, the stable phase of MAPbI1�xClx is sus-tainable only if the portion of chlorine replacement is less than 4%due to the large discrepancy in ionic radius between Cl� and I� [35].On the other hand, the mixed systemwith bromide (Br�) and iodide(I�) maintain stable phases without the limitation of the mixingratio because of less discrepancy in ionic radius [35,36]. The mixedelements in the ‘B’ site have been explored with lead (Pb) and tin(Sn). The ratio of Pb to Sn does not have the limitation with randomdistribution in the [MX6] octahedral, where M can be metal cationsite. The band structure and bandgap (Eg) changes according to theratio of Pb to Sn since the conduction bandminimum of the PVSKs ismainly originated from metal cation [37,38]. However, due to theinstability of Sn2þ, limited studies have been conducted [39]. Also,the mixed ‘A’ cations system has been investigated. The FAPbI3 hassmaller Eg and guarantees the longer carrier lifetime while its a-phase used for photovoltaic application is not stable under ambientcondition [9]. Partial substitution of FAwith smaller ‘A’ cations foundto stabilize the a-phase [19,40]. Now, the majority of high efficiencyPVSK solar cells are based on FA PVSKs with partial substitution ofthe FA cation with smaller MA, Cs, and Rb cations [13]. All materialsphases are mainly determined by the thermodynamics, changingaccording the temperature, pressure, and chemical potentials.Accordingly, the phase of the PVSK materials varies depending onthe given situation [41]. (Table 1) Of course, the properties of eachphase are completely different including crystal structure, electricalproperties and stability.

Table 1Crystallographic information of various halide PVSK materials. Reprinted with permissio

Phase Structure Lattice pa

MAPbl3 a Tetragonal a ¼ 6.311MAPbl3 b Tetragonal a ¼ 8.849MAPbl3 g Orthorhombic a ¼ 5.673MAPbCl3 a Cubic a ¼ 5.675MAPbCl3 b Tetragonal a ¼ 5.655MAPbCl3 g Orthorhombic a ¼ 5.673MAPbBr3 a Cubic a ¼ 5.901MAPbBr3 b Tetragonal a ¼ 8.322MAPbBr3 g Tetragonal a ¼ 5.894MAPbBr3 d Orthorhombic a ¼ 7.979MASnl3 a Tetragonal a ¼ 6.230MASnl3 b Tetragonal a ¼ 8.757FAPbl3 a Trigonal a ¼ 8.981FAPbl3 b Trigonal a ¼ 17.79FASnl3 a Orthorhombic a ¼ 6.328FASnl3 b Orthorhombic a ¼ 12.51

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The optoelectronic properties change depending on the crys-tallinity of the PVSK materials because the crystallinity determinesthe defect density of materials. The grain boundaries (GBs) play therole of crystal defects, which changes the material properties. Inthis section, we deal with the methods to grow single crystal andpolycrystalline PVSKs and their associated differences in properties.

Properties of the single crystal can provide a guideline for thedesign of materials and devices. For the single crystal halide PVSK,various growth techniques have been introduced including a top-seeded solution growth (TSSG), a hot-solution cooling-downgrowth, the inverse temperature method, and an anti-solvent va-por-assisted crystallization (Fig. 1) [12,25,42,43]. Regardless of themethod, they all follow the same principles where the key steps arenucleation and crystal growth. Controlling the nucleation sites andunderstanding the growth kinetics are the most important funda-mental knowledge. In addition, since all methods are based on thesolution process, the growth and dissolution happen at the sametime. In this case, controlling the growth rate is critical. Forexample, when the growth rate is too fast, the quality of the singlecrystal is not good enough because of the lower activation energyinduces more nucleation sites, whereas dissolution disturbs thegrowth of single crystal halide PVSKs when the growth rate is tooslow, resulting in poor crystal quality. The TSSG method utilizes asmall temperature gradient between the top and the bottom of thesolution and the temperature gradient induces convection toconvey the materials to the seed to enlarge the crystal [25]. Withthe beginning of MAPbI3, MAPbBr3 and MAPbCl3, even mixedhalide single crystals were also demonstrated [44]. The coolingdown method utilized the control of the amount of saturatedprecursors using Pb(CH3COOH)2$3H2O in hydroiodic acid (HI) tomix with MAI [12]. However, the growth rate is quite low in thiscase because the halide PVSK has inverse tendency of solubilityaccording to the temperature. That is, it has lower solubility in thesolvent at higher temperature unlike the common case. Due to thisreason, the inverse temperature method has been introduced touse the inverse tendency of solubility [45]. In this case, the growthrate becomes much faster. Because the antisolvent can work asanother source to control the chemical potential, the single crystalis grown by controlling the antisolvent. Specifically, the amount ofantisolvent helps to control the point where supersaturation initi-ates [46,47].

The optoelectronic properties of single crystal halide PVSKs varydepending on the characterization and growth methods. However,in general, the single crystal PVSKs do not have GBs and have fewerdefects, which leads to reduced recombination. Also, the absorptionedge and external quantum efficiency extends up to 850 nm

n from Ref. [41].

rameters (Å) Volume (Å)

5 b ¼ 6.3115 c ¼ 6.3161 251.6b ¼ 8.849 c ¼ 12.642 990b ¼ 5.628 c ¼ 11.182 959.6e e 182.2e c ¼ 5.630 180.1b ¼ 5.628 c ¼ 11.182 375e e 206.3e c ¼ 11.833 819.4

2 e c ¼ 5.8612 e

b ¼ 8.580 c ¼ 11.849 811.12 b ¼ 6.2302 c ¼ 6.2316 241.887 b ¼ 8.7577 c ¼ 12.429 953.27 b ¼ 8.9817 c ¼ 11.006 768.91 b ¼ 17.791 c¼ 10.091 2988.46 b ¼ 8.9554 c ¼ 8.9463 507.032 b ¼ 12.512 c ¼ 12.509 1959.2


Fig. 1. Schemes and images describing methods to grow PVSK single crystals. (a) Top seeded solution growth (b) A hot-solution cooling-down method (c) an anti-solvent vapor-assisted crystallization method. Reprinted with permission from Ref. [46].

J.-W. Lee et al. / Materials Today Energy xxx (2017) 1e12 3

whereas those from polycrystalline films are only around 800 nm.Thereby, Voc and fill factor (FF) can increase. Thus, more than 25%performance enhancement can be expected with other de-velopments for the devices of single crystals [46,47]. However,although it has benefits in terms of carrier movement, the currentsingle crystals only have bulky shape. In other words, it cannot beused for thin film electronics like solar cells and light emitting di-odes. Thus, another approach is required to use it as a material inthin film electronics.

PVSK solar cells utilize polycrystalline thin films grown by so-lution process. As a result, the effects of grain size and morphologyon photovoltaic performance have been investigated. There are tworepresentativeways tomake polycrystalline halide PVSKs; one-step


and two-step methods [18,48]. The two-step method consists oftwo sequential deposition steps. The first step is to deposit PbI2 filmand a second step is reacting the deposited PbI2 with MAI either byvapor evaporation, spin coating or dipping to convert it intoMAPbI3[17,49,50]. According to the thermodynamic regulation, the grainsize (Y) of the resulting PVSK film fabricated by the two-stepmethod can be determined by following relationship,

lnY ¼ 32ss13kT

�kTVmðlnX � lnC0ðTÞÞ2

�



where ss1 is the average surface tension, k is Boltzmann constant, Tis temperature of the reaction, X is the MAI concentration, C0 isequilibrium concentration of MAI [51]. As a result, the grain size (Y)of the MAPbI3 can be controlled by changing the reaction temper-ature or concentration of MAI (Fig. 2) [51,52]. The photovoltaicperformance of the PVSK solar cells were found to be stronglyaffected by the grain size of the film, which indicates the impor-tance of GBs in PVSK solar cells. Furthermore, in the two-stepmethod, the film quality including morphology and optoelectricalproperties is strongly affected by the quality of the PbI2 layer[51e53]. Due to the hygroscopic nature of PbI2, the control of thehumidity and temperature during the deposition of PbI2 is criticalto produce the superb device performance [8]. Q. Chen et al.demonstrated a vapor-assisted two-step process [50]. The vapor-assisted solution process uses the phase transformation from va-por to solid phase at a second step. Unlike the PbI2, the sublimationtemperature of organic components, such as MAI and FAI, is rela-tively low. Accordingly, PbI2 dissolved in solvent is firstly depositedvia spin coating and later the vaporized MAI or FAI components aredeposited on the top of the PbI2 layer. They reported reliable deviceperformance with large grain size. However, the phase trans-formation of vapor to solid phase accompanies the slow kineticsand, thus, the deposition rate of vaporized materials is much lowerthan solution process as usual.

In the one-step method, all the precursors for PVSK are dis-solved in a single solution. The one-step method has shown poorfilm quality due to the difficulty in controlling the reaction kineticand, thus it produced poor device performance [48]. In this regard,the anti-solvent method was introduced to control the reactionkinetics. As an antisolvent, toluene is widely used because it doesnot dissolve the precursor MAI and PbI2 [20,54]. Instead, thetoluene was involved into the reaction kinetics to induce the fastnucleation and crystallization of the film by controlling chemicalpotential. M. Xiao el al. tested 12 solvents working as the

Fig. 2. (a) Schematic showing two-step spin-coating process for deposition of CH3NH3PbI3theoretical expectation of the grain size (Y, solid line). (cef) Scanning electron microscopic (Sbars are 500 nm. Reprinted with permission from Refs. [48,51,52].


antisolvent, and chlorobenzenewas found to be the best in terms ofthe device performance [55]. Another approach is to utilize theintermediate phase. Lewis acidebase adduct approach has beenintroduced by Park et al. [15,16] Lewis acid PbI2 was reported toform adduct by reacting with oxygen, nitrogen, and sulfur bearinglone pair electrons. Dimethylsulfoxide and thiourea has been usedto form the adducts with PbI2, MAI and FAI [15]. The formation ofintermediate adducts enables control of crystallization kinetics toform highly uniform and crystalline PVSK films. While the coverageand grain size can be controlled in some extent, the formation of GBis inevitable, which motivated researchers to study the roles ofthe GB.

2.2. Photo-physical and computational works on GBs

As the PVSK solar cells have demonstrated unprecedentedlyrapid enhancement of PCE, the optoelectronic properties have beenstudied using photo-physical characterization tools to find theorigin of the superior photovoltaic performance. Charge carrierlifetime and diffusion length have been investigated using micro-scopic photoluminescence (PL) measurements with aid of transientabsorption spectroscopy, in which both electron and hole diffusionlengths were determined to be ~100 nm for MAPbI3 and ~1 mm forMAPbI3�xClx, which is much longer than typical solution processedorganic materials (~10 nm) [10,26]. The superior charge trans-porting ability was regarded as one of the main advantages oforganometal halide PVSK to show superb photovoltaic perfor-mance. While the polycrystalline thin films were characterized intheir initial stages, researchers have started to characterize theoptoelectronic properties of single crystal PVSK materials toexplore intrinsic properties of PVSK materials [12,25]. Dong et al.synthesized the 3 mm thick single crystal using supersaturatedMAPbI3 solution [25]. Notably, the MAPbI3 single crystal demon-strated much longer diffusion length than polycrystalline thin film,

(MAPbI3). (b) Experimentally measured grain size of the MAPbI3 (filled squares) andEM) images of MAPbI3 film with different concentration of CH3NH3I solution. The scale



in which the diffusion length was determined to be exceed 3 mmunder weak light (0.003 mWcm�2) and 175 mm under 1 sun(100 mW/cm�2) illumination, respectively. Shi et al. also charac-terized the MAPbI3 single crystals, in which diffusion length wasmeasured to be ~8 mm [12]. The measured diffusion length of thesingle crystal is more than one order of magnitude longer thanpolycrystalline MAPbI3 film, which implies characteristic GBs inpolycrystalline thin film comparing with single crystal counterparts probably have significant influence on charge transportingmechanism.

To understand the origin of superior charge carrier transportingability of the PVSK materials, first principle density-functionaltheory (DFT) calculations have been done to investigate thedominant intrinsic defects in the PVSK materials [11]. It was foundthat the dominant point defects in MAPbI3 PVSK are p-type VPb andn-type MAi, and both defects with low formation energies createonly shallow states due to strong Pb lone-pair s orbital and I porbital antibonding coupling and high iconicity of MAPbI3. The lowtransition energy level of the dominant defects with low formationenergy can be a possible origin of the superior intrinsic chargecarrier transporting properties of PVSK materials since the shallowdefect states do not contribute to ShockelyeReadeHall (SRH) non-radiative recombination loss, which is generally originated fromdeep gap states with higher transition energy [11,28,56].

The similar atomistic defects at GBs can induce the additionalstructural disorder such as a dangle bond and uncharacteristicwrong bond like cationecation and anioneanion bonds. Yin et al.investigated impact of such structural disorders at GBs on elec-tronic quality of PVSK materials [12]. Fig. 3 shows calculated den-sity of state (DOS) with different structural models at GBs.

P3(111)

andP

5(310) GBs were considered for MAPbI3 whileP

3(112) GB inCdTe was considered for comparison. Red and blue (in the webversion) circles in Fig. 3a and b indicate possible anioneanion andcationecation bonds, respectively. The calculated density of state(DOS) was illustrated in Fig. 3def. It is worthy to note that the

Fig. 3. The atomic structures of MAPbI3 (a) S3(111) GB, (b) S5(310) GB and (c) CdTe S3(112models and partial DOS of atoms in GB cores. All the DOS are normalized to per formula u


PbePbwrong bonds in bothP

3(111) andP

5(310) GBs do not formsubgap states, which was thought to be due to relatively largedistance between two Pb cations (3.84 and 5.17 Å for

P3(111) andP

5(310) GB, respectively) compared to ionic radius of Pb2þ

(1.33Å). On the other hand, both cations and anions inducedwrongbonds in CdTe (Fig. 3c and f) forms deep subgap states in CdTe,which strongly contribute to the non-radiative SRH recombination.The results explain why the solution processed polycrystallinePVSK film can have such an excellent charge transporting proper-ties and resulting high open-circuit voltage (VOC) comparable tosingle-crystal GaAs thin film. However, although the wrong bondsin GBs of PVSK do not form deep trap states, they are still found toform shallow defect states. As can be seen in Fig. 3b and e, the IeIwrong bond in

P5(310) GB forms shallow defects states close to

valance band maximum (VBM). Although this shallow state is notas critical as deep levels as in CdTe, it possibly increase the effectivemass of holes or trap the holes to decrease the hole mobility orinduce non-radiative SRH recombination [12]. Long et al. alsoobserved accelerated non-radiative recombination at GBs throughtime domain atomistic simulation via nonadiabatic (NA) moleculardynamics combined with time-domain density functional theory[57]. While GBs were confirmed to not forms deep gap states, theyconcluded that GBs reduce the Eg and enhance the electron phononcoupling to accelerate the non-radiative charge recombination loss.

With expected detrimental effects of GBs on photovoltaic per-formance suggested from computational works, the presence ofsuch defect states have confirmed from experimental observation.Duan et al. observed presence of the defect states ~0.16 eV abovethe valance band using admittance spectroscopy [58]. The defecttrapping and resulting non-radiative recombination have beenobserved from microscopic PL measurement with much shorterdecay time constant (<10 ns) than radiative recombination at bulkPVSK (>100 ns) [59]. However, the microscopic characterizationtools could not specify the origin of the defect states and resultingnon-radiative recombination. Confocal fluorescence microscopy

) GB after structural relaxations. (def) Corresponding total density of state (DOS) of GBnit. Reprinted with permission from Ref. [12].



measurement was suggested as a powerful tool to investigatespatially resolved charge carrier recombination dynamics. de Qui-lettes et al. first incorporates the confocal PL to study local PL decaydynamic of the polycrystalline PVSK thin film [22]. Regardless oflong microscopic PL lifetime exceeding 1 ms, the steady-state PLintensity showed significant heterogeneity across the PVSK film.More importantly, closer inspection of local PL unraveled thatsteady-state PL intensity at GBs was lower than that of grain in-teriors, which was found to be due to much shorter charge carrierlifetime at GBs. Considering the local PL intensity and lifetime, itwas concluded that the non-radiative trap-assisted recombinationdue to defects is more severe at GBs, which is correlated withcomputational works mentioned above. More recently, Ham et al.also observed accelerated charge carrier recombination at GBs ofMAPbI3 PVSK films prepared from both stoichiometric and non-stoichiometric precursor solutions (Fig. 4) [60]. Statistical distri-bution of average charge PL lifetime (tavg) clearly indicates that thetavg is relatively longer at grain interiors compared to GBs regard-less of the composition of the precursor solution. The shorter PLlifetime was found to be due to contribution of trap-assisted non-radiative recombination, which is more dominant at GBs of MAPbI3compared to grain interiors.

However, there are still contradictory reports regarding impactof GBs on charge carrier lifetime. Yang et al. performed confocalfluorescence-lifetime imaging of MAPbI3 grown from non-stoichiometric precursor solution [61]. They observed comparablePL lifetime at grains interiors and GBs although the intensity of PL isweaker at GBs compared to grain interiors. They observed that therecombination in these film primarily occur in non-GB regions suchas grain surface or grain interiors. A recent study using a transientreflection spectroscopy by same group also pointed out that thecharge carrier lifetime is mainly limited by surface recombination

Fig. 4. Confocal photoluminescence images of PVSK films prepared from (a) stoichiometric (is 500 nm. (c,d) Two-dimensional scatter plots of PL intensity and average lifetime for (c) scurves (excitation at 420 nm, 250 kHz with fluence of ~0.8 mJ/cm2). PL decay curves at graincurves at grain (blue square; ,) and GB (green diamond; ◊) in nonst-PVSK film are showncolour in this figure legend, the reader is referred to the web version of this article.)


[62]. However, considering the PVSK films used for the character-ization were prepared from non-stoichiometric solution (MAIexcess), the chemical environment at GBs probably changed, whichwill affect the electronic state and resulting charge carrier behavior.Further systematic photo-physical and computational studies onGBs with different chemical environments might be necessary tofully understand the role of GBs in PVSK solar cells.

2.3. AFM and KPFM studies on GBs

In analogous to the local PLmeasurement, the Kelvin probe forcemicroscopy (KPFM) and atomic force microscopy (AFM) are effi-cient tools to investigate the local properties of the halide PVSK(Fig. 5a and b). Themain function of KPFM is tomeasure the contactpotential difference from the vacuum level and the surface photo-voltage. The KPFM is taken at a noncontact mode. In noncontactmode, a conducting cantilever is scanned over the surface of sam-ples. The cantilever keeps a constant height to measure the contactpotential difference between samples and the cantilever which isregarded as a reference electrode. Thus, work function of thesamples can be characterized. On the other hand, local photocur-rent of the halide PVSK materials can be measured using con-ducting AFM (c-AFM) mode. Whereas AFM is utilized to measurethe topography of the surface, the c-AFM can be used to measurethe electric current flow between a conducting cantilever andsamples. The c-AFM is operated in contact mode. The cantilever isscanned under a constant voltage to get current information.

Structural defects at GBs can induce change in electronic bandstructure. For example, GBs in polycrystalline Si thin film createsdeep energy states between the conduction band and the valenceband, which works as a recombination site [63,64]. Whereas, GBs inchalcogenide compounds such as CZTS, CZTSSe, CIGS, have

st) solution and (b) non-stoichiometric (nonst) solution (MAI:PbI2 ¼ 1.06:1.0). Scale bart-PVSK and (d) nonst-PVSK MAPbI3 films. (e,f) Representative time-resolved PL decay(red circle; B) and GB (yellow triangle; △) in st-PVSK films are shown in (e). PL decayin (f). Reprinted with permission from Ref. [60]. (For interpretation of the references to


Fig. 5. (a) Kelvin probe force microscopy (KPFM) and (b) conducting atomic force microscopy (c-AFM) setup for characterization of PVSK layer. (c) KPFM and (d) c-AFM images of thePVSK layer. Inset of (c) and (d) show topological images. Reprinted with permission from Refs. [23,29].


beneficial effects in terms of minority carrier collections [65]. It isreported that GBs in the compounds do not produce deep stateswithin the Eg. Instead, the GBs are electrically charged and inducebuilt-in potential. As a result, photo-generated electrons and holenear the GBs can be easily separated by the built-in potential, whichenhances photocurrent. In the same manner, Yun et al. found outthat the GBs in PVSK can help the charge separation and collectionby observing the KPFM and AFM results in PVSK films [23]. Theyclaimed that the potential barrier along the GB was observed underdark conditions whereas higher photovoltage was obtained alongthe GB compared with the grain interior under the illumination(Fig. 5c). Thus, the higher current collection near GBs is observed intheir work, which supports the beneficial role of GBs in terms ofcharge collection. Q. Chen et al. also studied contact potential dif-ference between near and at GBs [66]. They employed scanningkelvin probe microscopy (SKPM) to examine the surface and GBsand concluded that there is band bending near the GBs. Releasingthe organic species to form PbI2 changes the grain to GB bendbending from downward to upward, preventing the chargerecombination at the interface between PVSK and a hole trans-portingmaterial. Son et al. reported that the self-formed GB healinglayer using 6 mol% excess CH3NH3I [29]. More efficient current flowthrough the GBs was confirmed by c-AFM, which can be furtherenhanced by the 6 mol% excess CH3NH3I (Fig. 5d). Recently, theheterogeneity of halide PVSK materials was reported by Leblebiciet al. using c-AFM [24]. Local characterization by the c-AFM allowedthem to investigate the microscopic and nanoscopic photovoltaicperformance of the PVSK layers. The heterogeneity is not onlypresented between the intra-grain and GBs but also evenwithin theintra-grain. They insist that difference in defect density withdifferent crystal facet contributes to the heterogeneity in localphotovoltaic performance, which eventually influences themacroscale solar cell performance.


2.4. Hysteresis and ion migration

While characterizing photovoltaic performance of PVSK solarcells, researchers had observed the irregular currentevoltage (IeV)characteristics (i.e. the IeV curve of PVSK solar cells variesdepending on the scan direction and the scan speed) at the earlystage, which was defined as IeV hysteresis. Three possible mech-anisms of the hysteresis phenomenon have been proposed: 1)trapping and detrapping effect on charge carriers. 2) polarizationdue to the ferroelectric properties. 3) ion migration [67,68]. Sup-ported by many direct and indirect evidences, such as the activeenergy of ion movements and the response time scale of the hys-teresis, it has been believed that ionmigration plays a crucial role inthe hysteresis behavior of PVSK solar cells [29,68e71].

Under external electric field, ions are driven to the specific di-rection and accumulate at the interfaces between the PVSK and thecharge transporting layers to form the doped regions (as illustratedat Fig. 6) [69]. The accumulated ions can induce an internal electricfield near the interface (space charge), screening the appliedvoltage, which induce the IeV hysteresis [29,68e70]. The evidenceof ion migration in the PVSK device has been distinctly reportedthrough the observation of the field-switchable photovoltaic effect,in which the photocurrent direction can be switched repeatedly byapplying a small electric field [69].

The existence of defects in the condensedmaterials provides themedium for ion migration. The possibility of ion migration highlydepends on their activation energy (EA). There are several factorscontributing to the activation energy, such as crystal structure, ionradius and valence state of ions. The ion migration rate rm can beestimated by Arrhenius relation: rmfexp

�� EAkBT

�, where kBT is the

thermal energy with kB for the Boltzmann constant and T for ab-solute temperature. A series of experimental studies combinedwith the first principles calculations has clearly revealed the


Fig. 6. Illustration of induced ion drift within PVSK device by (a) positive and (b) negative poling. The p- and n-doped regions arise from the ion accumulation at interface. (c,d)Three ionic transport mechanisms in MAPbI3 PVSKs. (c) Pb and I vacancies. (d) MA vacancy. Reprinted with permission from Refs. [69,72].


possible mobile ion species and migration paths in PVSK com-pounds [29,72e76]. Taking MAPbI3 PVSK as an example, the mostlikely mobile ion is considered to be iodide ions. Eames and hiscoworkers identified the ionic transport mechanisms involvingvacancy hoping between neighboring positions (as shown in Fig. 7),in which Pb2þ ions move along the diagonal <110> direction, I�

Fig. 7. The illustration of the ion migration induced by different mechanisms. (a) Schottkysolved impurities, (e) soften lattice by lamination, (f) GBs and (g) strain induced defect. Re


ions migrate through the octahedron edge, andMAþ ions jump intothe nearest vacant cages [72]. Based on the density functionaltheory methods, they found the facial vacancy-assisted migrationof I� ions with activation energy of 0.58 eV that is lower than theone of MAþ (0.84 eV) and Pb2þ (2.31 eV). By the comparison withthe result of the temperature- and time-dependent kinetic

defects, (b) Frenkel defects, (c) Lattice distortion by accumulated charges and (d) dis-printed with permission from Ref. [29].



measurements, their calculated activation energy for iodide ionsmigration fits well for the formation of hysteresis (0.60e0.68 eV)[72]. Another theoretical computations performed by Azpiroz et al.also shows the identical tendency with regard to the activationenergy of ion migration via vacancy hopping in the similar migra-tion paths. Conversely, they suggest that MAþ ions migration withEA of 0.5 eV is responsible for the slow response in PVSK solar cells.According to their prediction, the iodide ion migration with thesmall EA (0.08 eV) can only occur within a very short time (<1 ms)that is too fast to explain the slow response of hysteresis effect inPVSK (usually in the order between few seconds to minutes) [73].

Besides through the bulk point defects (i.e. Schottky and Frankeldefects) inside the grain interior, ions can also migrate along theGBs preserving high density of defects. With more open space andlack of chemical bonding at GBs, the activation energy of ionmigration is expected to become lower than the one in the graininterior. In addition to point defects and GBs, soften lattice andlattice distortion can also be the causes of ion migration. Thepossible pathways for the ionmigration are illustrated in Fig. 7 [29].

Recently, several experimental observations have indicated thatGBs might play a critical role while serving as an ion migrationchannel. In Xiao et al.'s work, the grain size dependence of theswitchable photovoltaic effect in symmetrical lateral PVSK deviceswas investigated. They found that it is easier to polarize the filmswith more GBs, and attribute it to the profound vacancy concen-tration [69]. In addition, Shao et al. also observed the grain sizedependent ion migration by analyzing the measurement ofconducting-AFM, SEM and EDS. They revealed that ion migration inpolycrystalline PVSKs dominates through GBs evidenced by distinctelement redistribution along GBs after electrical poling [30]. Byusing KPFM and AFM, Yun et al. confirmed that GBs can act asshortcuts for ion migration [77]. The ions move faster at the GBscompared to the grain interior owing to higher ionic diffusivity andhigher segregation of ions at the GBs. The result was correlatedwith the report by Xing et al.; a single crystal PVSK possess at leasttwo-fold higher activation energy for ion movement than the onein the polycrystalline thin film, which probably due to absence ofGBs [78].

The GBs were found to have important role in stability of thePVSK solar cells [79e81]. Redistribution of the ions under operationpossibly causes the degradation of device performance. Domanski

Fig. 8. Schematics showing proposed mechanisms for PbI2 passivatio


et al. reported the loss of the performance bymigration of cations inPVSK layer although it is observed to be reversible upon storageunder dark [81]. Furthermore, the ingression of moisture occurpreferentially from GBs, and as a results, the films with more GBsdemonstrated poor moisture stability, which also observed interms of device stability [79,80]. Therefore, the growth of perov-skite films with larger grain is thought to be better to improve thestability of the device.

2.5. Passivation of GBs

It has been widely shown that varying the stoichiometry of theprecursor solutions and subsequent solid-state films can affect theoptoelectronic properties of PVSK thin films. For the one-stepprecursor deposition technique, this is achieved in a controlledmanner by intentionally altering the molar ratio off a one-to-onestoichiometry by adding additional methylammonium halide(MAX/FAX) or lead halide (PbX2) to add excess organic or leadcontent, respectively. For two-step deposition techniques, it is moredifficult to precisely control stoichiometry, where typically excessPbI2 is formed at GBs during prolonged annealing due to thermaldecomposition of species. The effect of thermally induced PbI2 onphotovoltaic performance was first reported by Chen et al. [66]Their results demonstrated that thermally induced PbI2 formed atGBs could provide beneficial passivation effects by suppressingrecombination between the selective contacts and PVSK, thusimproving the photovoltaic performance as shown in Fig. 8.

Excess MAI has also been shown to lead to improved carrierlifetimes and performances in PVSK thin films [29,60]. It was foundthat excessMAIwas able to provide beneficial passivating effects onGBs to achieve higher carrier lifetimes and photoluminescence in-tensities as was evidenced through spatially and temporallyresolved photoluminescencemeasurements, ultimately resulting inhigher device performances [60]. The improvements were attrib-uted to a smaller initial decay component resulting from passiv-ation effects near GBs through excess iodine that served to suppressnon-radiative recombination as well as improve the quality of theindividual grains in terms of morphology and Schottky disorder.Moreover, excess MAI at GBs has shown to provide a ‘conductivepathway’ due to their formation and passivating effects at GBs, asshown through conducting atomic force microscopy (c-AFM)

n in PVSK solar cells. Reprinted with permission from Ref. [66].



measurements in previous section (Fig. 5d). Ultimately, these de-vices were able to achieve efficiencies exceeding 20% [29].

In terms of halide substitution, small amounts of chlorineincorporation have been widely shown to improve performancesand carrier lifetimes of PVSK thin films. Yin et al. found throughcomputational studies that Cl and O can spontaneously segregatewithin GBs and passivate defect/trap states [12], thus explaininghow such small amounts of Cl incorporation can have significanteffects on carrier lifetime and performance. In addition, chlorineinclusion has shown to assist in film formation effects, thusimproving the grain quality and mitigating defect-induced non-radiative recombination at GBs. Bromine has been another halidethat has been used to partially substitute for iodine to improveperformances [82,83]. Bromine partial substitution has shown tohave effects on GB passivation and surface defect reduction. Forinstance, Yang and company used a MABr surface treatment toconvert MAPbI3 thin films to higher qualities through an Ostwald-ripening process with a IeBr anion exchange [83]. This led to notonly higher performances but devices with higher stability as well.Ultimately, varying halide species can result in enhanced crystal-linity by changing the nature of the film growth, due to changes insolubility of precursors, changing reaction pathways, etc, and alsoprovide passivation effects at GBs when in excess. Therefore, off-stoichiometry presents one viable technique to improve the opto-electronic properties of PVSK films towards higher performances.

Aside from adjusting the stoichiometry of the PVSK precursorcomponents, additives have been used as GB engineering tech-niques to provide further control over the nucleation and growthprocesses of PVSK thin films, and serve as passivating species toimprove optoelectronic properties. Incorporation of a small amountof guanidinium halides has shown to improve film quality for two-step techniques (Fig. 9) [84]. The additives had effects on film for-mation, where the morphology was improved with less prominentGBs (Fig. 9a and b). In addition, a three-fold increase in

Fig. 9. Scanning electron microscopic (SEM) images of PVSK film (a) without and (b) withspectra of the PVSK films. Reprinted with permission from Ref. [84].


photoluminescence intensity and order of magnitude increase incarrier lifetime was observed for PVSK films with the guanidiniumadditives (Fig. 9c and d). This led to higher device efficiency with aVoc over 1.1 V for planar two-step devices. The improvements wereattributed to reduced defect densities and GB passivation effects viathe guanidinium-based additives [84]. Lead thiocyanate has alsobeen used as an additive to improve the properties of the PVSKfilms, where HSCN and CH3NH3 gases were shown to evolve fromreactions during film formation that resulted in the formation ofexcess PbI2 at GBs, providing passivating effects [85]. Hypophos-phorous acid (HA) as an additive has shown to improve grainquality and performance, which the authors attribute to reducingthe defect density at GBs and surfaces during film growth [40].KPFM measurements showed a reduction in surface potential of210 mV that changes the band bending, the magnitude of whichindicates the density of surface trap states, while fluorescencemicroscopy analysis showed enhanced fluorescence and lifetimesdue to HA addition.

Since PVSKs are ionic compounds, the constitute ions can formchemical interactions with added chemicals. For example, under-coordinated Pb2þ ions at GBs and surface of grain can be passivatedby interacting with lone pair electrons in pyridine [86]. Further-more, polymer-templated growth of PVSK films has recently beendemonstrated by Bi et al., where GBs grew perpendicular to thesubstrate under the influence of the PMMA additive, which effec-tively reduced the surface area of the boundaries [87]. Thiscontrolled growth using PMMA led to long photoluminescencelifetimes and a certified power conversion efficiency over 21%. Thecommonality behind additive-based GB engineering is that addi-tives can be used to control crystallization processes and enhancefilm quality, while also providing passivating effects via remnantspecies within the film. Additives are becoming more widelystudied in more detail and represents a powerful technique toenhance optoelectronic quality of PVSK thin films.

guadinium iodide additive. (c) Steady-state and (d) time-resolved photoluminescence



3. Conclusion

Investigation of underlying microscopic properties of the PVSKsbecomes more important to understand the origin of their bulkproperties. Although it is still unclear whether the GBs are bene-ficial or detrimental on photovoltaic performance and stability ofthe PVSK solar cells, they undoubtedly play important roles inoptoelectronic and physical processes in the PVSK solar cells.Further works will be required to fully understand the role of GBs.We expect the understanding of microscopic properties will pro-vide the important insight to improve the performance and sta-bility of the PVSK solar cell towards commercialization.

Acknowledgements

This work is supported by Air Force Office of Scientific Research(AFOSR, Grant No. FA9550e15-1e0610 and FA9550-15-1-0333),Office of Naval Research (ONR, Grant No. N00014-04-1-0434), andNational Science Foundation (NSF, Grant No. DMR-1210893 andECCS-EPMD-1509955).

References

[1] A. Kojima, K. Teshima, Y. Shirai, T. Miyasaka, Organometal halide perovskitesas visible-light sensitizers for photovoltaic cells, J. Am. Chem. Soc. 131 (2009)6050e6051.

[2] J.-H. Im, C.-R. Lee, J.-W. Lee, S.-W. Park, N.-G. Park, 6.5% efficient perovskitequantum-dot-sensitized solar cell, Nanoscale 3 (2011) 4088e4093.

[3] H.-S. Kim, C.-R. Lee, J.-H. Im, K.-B. Lee, T. Moehl, A. Marchioro, S.-J. Moon,R. Humphry-Baker, J.-H. Yum, J.E. Moser, Lead iodide perovskite sensitized all-solid-state submicron thin film mesoscopic solar cell with efficiencyexceeding 9%, Sci. Rep. 2 (2012) 591.

[4] M.M. Lee, J. Teuscher, T. Miyasaka, T.N. Murakami, H.J. Snaith, Efficient hybridsolar cells based on meso-superstructured organometal halide perovskites,Science 338 (2012) 643e647.

[5] National renewable energy laboratory, Best research-cell efficiency chart,https://www.nrel.gov/pv/assets/images/efficiency-chart.png.

[6] J.-W. Lee, T.-Y. Lee, P.J. Yoo, M. Gr€atzel, S. Mhaisalkar, N.-G. Park, Rutile TiO2-based perovskite solar cells, J. Mater. Chem. A 2 (2014) 9251e9259.

[7] M. Yang, Z. Li, M.O. Reese, O.G. Reid, D.H. Kim, S. Siol, T.R. Klein, Y. Yan,J.J. Berry, M.F.A.M. van Hest, K. Zhu, Perovskite ink with wide processingwindow for scalable high-efficiency solar cells, Nat. Energy 2 (2017) 17038.

[8] H.-S. Ko, J.-W. Lee, N.-G. Park, 15.76% efficiency perovskite solar cells preparedunder high relative humidity: importance of PbI2 morphology in two-stepdeposition of CH3NH3PbI3, J. Mater. Chem. A 3 (2015) 8808e8815.

[9] J.W. Lee, D.J. Seol, A.N. Cho, N.G. Park, High-efficiency perovskite solar cellsbased on the black polymorph of HC(NH2)2PbI3, Adv. Mater 26 (2014)4991e4998.

[10] G. Xing, N. Mathews, S. Sun, S.S. Lim, Y.M. Lam, M. Gr€atzel, S. Mhaisalkar,T.C. Sum, Long-range balanced electron-and hole-transport lengths inorganic-inorganic CH3NH3PbI3, Science 342 (2013) 344e347.

[11] W.-J. Yin, T. Shi, Y. Yan, Unusual defect physics in CH3NH3PbI3 perovskite solarcell absorber, Appl. Phys. Lett. 104 (2014) 063903.

[12] W.J. Yin, H. Chen, T. Shi, S.H. Wei, Y. Yan, Origin of high electronic quality instructurally disordered CH3NH3PbI3 and the passivation effect of Cl and O atgrain boundaries, Adv. Electron. Mater. 1 (2015) 1500044.

[13] M. Saliba, T. Matsui, K. Domanski, J.-Y. Seo, A. Ummadisingu,S.M. Zakeeruddin, J.-P. Correa-Baena, W.R. Tress, A. Abate, A. Hagfeldt,Incorporation of rubidium cations into perovskite solar cells improvesphotovoltaic performance, Science 354 (2016) 206e209.

[14] W.E. Sha, X. Ren, L. Chen, W.C. Choy, The efficiency limit of CH3NH3PbI3perovskite solar cells, Appl. Phys. Lett. 106 (2015) 221104.

[15] J.-W. Lee, H.-S. Kim, N.-G. Park, Lewis acidebase adduct approach for highefficiency perovskite solar cells, Acc. Chem. Res. 49 (2016) 311e319.

[16] N. Ahn, D.-Y. Son, I.-H. Jang, S.M. Kang, M. Choi, N.-G. Park, Highly repro-ducible perovskite solar cells with average efficiency of 18.3% and best effi-ciency of 19.7% fabricated via Lewis base adduct of lead (II) iodide, J. Am.Chem. Soc. 137 (2015) 8696e8699.

[17] J. Burschka, N. Pellet, S.-J. Moon, R. Humphry-Baker, P. Gao, M.K. Nazeeruddin,M. Gr€atzel, Sequential deposition as a route to high-performance perovskite-sensitized solar cells, Nature 499 (2013) 316e319.

[18] J.-W. Lee, N.-G. Park, Two-step deposition method for high-efficiency perov-skite solar cells, MRS Bull. 40 (2015) 654e659.

[19] J.W. Lee, D.H. Kim, H.S. Kim, S.W. Seo, S.M. Cho, N.G. Park, Formamidinium andcesium hybridization for photo-and moisture-stable perovskite solar cell, Adv.Energy Mater. 5 (2015) 1501310.


[20] N.J. Jeon, J.H. Noh, W.S. Yang, Y.C. Kim, S. Ryu, J. Seo, S.I. Seok, Compositionalengineering of perovskite materials for high-performance solar cells, Nature517 (2015) 476e480.

[21] C.C. Stoumpos, C.M.M. Soe, H. Tsai, W. Nie, J.-C. Blancon, D.H. Cao, F. Liu,B. Traor�e, C. Katan, J. Even, A.D. Mohite, M.G. Kanatzidis, High members of the2D Ruddlesden-popper halide perovskites: synthesis, optical properties, andsolar cells of (CH3(CH2)3NH3)2(CH3NH3)4Pb5I16, Chem 2 (2017) 427e440.

[22] D.W. de Quilettes, S.M. Vorpahl, S.D. Stranks, H. Nagaoka, G.E. Eperon,M.E. Ziffer, H.J. Snaith, D.S. Ginger, Impact of microstructure on local carrierlifetime in perovskite solar cells, Science 348 (2015) 683e686.

[23] J.S. Yun, A. Ho-Baillie, S. Huang, S.H. Woo, Y. Heo, J. Seidel, F. Huang, Y.-B. Cheng, M.A. Green, Benefit of grain boundaries in organiceinorganic halideplanar perovskite solar cells, J. Phys. Chem. Lett. 6 (2015) 875e880.

[24] S.Y. Leblebici, L. Leppert, Y. Li, S.E. Reyes-Lillo, S. Wickenburg, E. Wong, J. Lee,M. Melli, D. Ziegler, D.K. Angell, Facet-dependent photovoltaic efficiencyvariations in single grains of hybrid halide perovskite, Nat. Energy 1 (2016)16093.

[25] Q. Dong, Y. Fang, Y. Shao, P. Mulligan, J. Qiu, L. Cao, J. Huang, Electron-holediffusion lengths >175 mm in solution-grown CH3NH3PbI3 single crystals,Science 347 (2015) 967e970.

[26] S.D. Stranks, G.E. Eperon, G. Grancini, C. Menelaou, M.J. Alcocer, T. Leijtens,L.M. Herz, A. Petrozza, H.J. Snaith, Electron-hole diffusion lengths exceeding 1micrometer in an organometal trihalide perovskite absorber, Science 342(2013) 341e344.

[27] D.-Y. Son, J.-W. Lee, Y.J. Choi, I.-H. Jang, S. Lee, P.J. Yoo, H. Shin, N. Ahn, M. Choi,D. Kim, Self-formed grain boundary healing layer for highly efficientCH3NH3PbI3 perovskite solar cells, Nat. Energy 1 (2016) 16081.

[28] J. Kim, S.-H. Lee, J.H. Lee, K.-H. Hong, The role of intrinsic defects in methyl-ammonium lead iodide perovskite, J. Phys. Chem. Lett. 5 (2014) 1312e1317.

[29] Y. Yuan, J. Huang, Ion migration in organometal trihalide perovskite and itsimpact on photovoltaic efficiency and stability, Acc. Chem. Res. 49 (2016)286e293.

[30] Y. Shao, Y. Fang, T. Li, Q. Wang, Q. Dong, Y. Deng, Y. Yuan, H. Wei, M. Wang,A. Gruverman, Grain boundary dominated ion migration in polycrystallineorganiceinorganic halide perovskite films, Energy Environ. Sci. 9 (2016)1752e1759.

[31] C. Li, X. Lu, W. Ding, L. Feng, Y. Gao, Z. Guo, Formability of ABX3 (X¼F, Cl, Br, I)halide perovskites, Acta Crystallogr. Sect. B Struct. Sci. 64 (2008) 702e707.

[32] C.C. Stoumpos, C.D. Malliakas, M.G. Kanatzidis, Semiconducting tin and leadiodide perovskites with organic cations: phase transitions, high mobilities,and near-infrared photoluminescent properties, Inorg. Chem. 52 (2013)9019e9038.

[33] D. Liu, T.L. Kelly, Perovskite solar cells with a planar heterojunction structureprepared using room-temperature solution processing techniques, Nat. Pho-tonics 8 (2014) 133e138.

[34] Q. Chen, H. Zhou, Y. Fang, A.Z. Stieg, T.-B. Song, H.-H. Wang, X. Xu, Y. Liu, S. Lu,J. You, The optoelectronic role of chlorine in CH3NH3PbI3 (Cl)-based perovskitesolar cells, Nat. Commun. 6 (2015) 7269.

[35] Y. Wu, A. Islam, X. Yang, C. Qin, J. Liu, K. Zhang, W. Peng, L. Han, Retarding thecrystallization of PbI2 for highly reproducible planar-structured perovskitesolar cells via sequential deposition, Energy Environ. Sci. 7 (2014) 2934e2938.

[36] F. Hao, C.C. Stoumpos, D.H. Cao, R.P. Chang,M.G. Kanatzidis, Lead-free solid-stateorganic-inorganic halideperovskite solar cells, Nat. Photonics 8 (2014) 489e494.

[37] F. Hao, C.C. Stoumpos, R.P. Chang, M.G. Kanatzidis, Anomalous band gapbehavior in mixed Sn and Pb perovskites enables broadening of absorptionspectrum in solar cells, J. Am. Chem. Soc. 136 (2014) 8094e8099.

[38] W.-J. Yin, J.-H. Yang, J. Kang, Y. Yan, S.-H. Wei, Halide perovskite materials forsolar cells: a theoretical review, J. Mater. Chem. A 3 (2015) 8926e8942.

[39] N.K. Noel, S.D. Stranks, A. Abate, C. Wehrenfennig, S. Guarnera, A.-A. Haghighirad, A. Sadhanala, G.E. Eperon, S.K. Pathak, M.B. Johnston, Lead-free organiceinorganic tin halide perovskites for photovoltaic applications,Energy Environ. Sci. 7 (2014) 3061e3068.

[40] W. Zhang, S. Pathak, N. Sakai, T. Stergiopoulos, P.K. Nayak, N.K. Noel,A.A. Haghighirad, V.M. Burlakov, A. Sadhanala, W. Li, Enhanced optoelectronicquality of perovskite thin films with hypophosphorous acid for planar het-erojunction solar cells, Nat. Commun. 6 (2015) 10030.

[41] S.-H. Bae, H. Zhao, Y.-T. Hsieh, L. Zuo, N. De Marco, Y.S. Rim, G. Li, Y. Yang,Printable solar cells from advanced solution-processible materials, Chem 1(2016) 197e219.

[42] Y. Dang, Y. Liu, Y. Sun, D. Yuan, X. Liu, W. Lu, G. Liu, H. Xia, X. Tao, Bulk crystalgrowth of hybrid perovskite material CH3NH3PbI3, CrystEngComm 17 (2015)665e670.

[43] Q. Han, S.H. Bae, P. Sun, Y.T. Hsieh, Y.M. Yang, Y.S. Rim, H. Zhao, Q. Chen,W. Shi, G. Li, Single crystal formamidinium lead iodide (FAPbI3): insight intothe structural, optical, and electrical properties, Adv. Mater. (2016)2253e2258.

[44] Y. Dang, Y. Zhou, X. Liu, D. Ju, S. Xia, H. Xia, X. Tao, formation of hybridperovskite tin iodide single crystals by top-seeded solution growth, Angew.Chem. 128 (2016) 3508e3511.

[45] M.I. Saidaminov, A.L. Abdelhady, B. Murali, E. Alarousu, V.M. Burlakov,W. Peng, I. Dursun, L. Wang, Y. He, G. Maculan, High-quality bulk hybridperovskite single crystals within minutes by inverse temperature crystalli-zation, Nat. Commun. 6 (2015) 7586.

[46] N.-G. Park, Crystal growth engineering for high efficiency perovskite solarcells, CrystEngComm 18 (2016) 5977e5985.


http://refhub.elsevier.com/S2468-6069(17)30123-5/sref1















https://www.nrel.gov/pv/assets/images/efficiency-chart.png

















































































































































































































[47] J. Huang, Y. Shao, Q. Dong, Organometal trihalide perovskite single crystals: anext wave of materials for 25% efficiency photovoltaics and applicationsbeyond? J. Phys. Chem. Lett. 6 (2015) 3218e3227.

[48] J.-H. Im, H.-S. Kim, N.-G. Park, Morphology-photovoltaic property correlationin perovskite solar cells: one-step versus two-step deposition of CH3NH3PbI3,APL Mater. 2 (2014) 081510.

[49] K. Liang, D.B. Mitzi, M.T. Prikas, Synthesis and characterization of organ-ic�inorganic perovskite thin films prepared using a versatile two-step dip-ping technique, Chem. Mater. 10 (1998) 403e411.

[50] Q. Chen, H. Zhou, Z. Hong, S. Luo, H.-S. Duan, H.-H. Wang, Y. Liu, G. Li, Y. Yang,Planar heterojunction perovskite solar cells via vapor-assisted solution pro-cess, J. Am. Chem. Soc. 136 (2013) 622e625.

[51] N. Ahn, S.M. Kang, J.-W. Lee, M. Choi, N.-G. Park, Thermodynamic regulation ofCH3NH3PbI3 crystal growth and its effect on photovoltaic performance ofperovskite solar cells, J. Mater. Chem. A 3 (2015) 19901e19906.

[52] J.-H. Im, I.-H. Jang, N. Pellet, M. Gr€atzel, N.-G. Park, Growth of CH3NH3PbI3cuboids with controlled size for high-efficiency perovskite solar cells, Nat.Nanotechnol. 9 (2014) 927e932.

[53] D.J. Seol, J.W. Lee, N.G. Park, On the role of interfaces in planar-structuredHC(NH2)2PbI3 perovskite solar cells, ChemSusChem 8 (2015) 2414e2419.

[54] N.J. Jeon, J.H. Noh, Y.C. Kim, W.S. Yang, S. Ryu, S.I. Seok, Solvent engineering forhigh-performance inorganiceorganic hybrid perovskite solar cells, Nat. Mater.13 (2014) 897e903.

[55] M. Xiao, F. Huang, W. Huang, Y. Dkhissi, Y. Zhu, J. Etheridge, A. Gray-Weale,U. Bach, Y.B. Cheng, L. Spiccia, A fast deposition-crystallization procedure forhighly efficient lead iodide perovskite thin-film solar cells, Angew. Chem. 126(2014) 10056e10061.

[56] W.J. Yin, T. Shi, Y. Yan, Unique properties of halide perovskites as possibleorigins of the superior solar cell performance, Adv. Mater. 26 (2014)4653e4658.

[57] R. Long, J. Liu, O.V. Prezhdo, Unravelling the effects of grain boundary andchemical doping on electronehole recombination in CH3NH3PbI3 perovskiteby time-domain atomistic simulation, J. Am. Chem. Soc. 138 (2016)3884e3890.

[58] H.-S. Duan, H. Zhou, Q. Chen, P. Sun, S. Luo, T.-B. Song, B. Bob, Y. Yang, Theidentification and characterization of defect states in hybrid organ-iceinorganic perovskite photovoltaics, Phys. Chem. Chem. Phys. 17 (2015)112e116.

[59] X. Wen, Y. Feng, S. Huang, F. Huang, Y.-B. Cheng, M. Green, A. Ho-Baillie,Defect trapping states and charge carrier recombination in organiceinorganichalide perovskites, J. Mater. Chem. C 4 (2016) 793e800.

[60] S. Ham, Y.J. Choi, J.-W. Lee, N.-G. Park, D. Kim, Impact of excess CH3NH3I onfree carrier dynamics in high-performance non-stoichiometric perovskites,J. Phys. Chem. C 121 (2017) 3143e3148.

[61] M. Yang, Y. Zeng, Z. Li, D.H. Kim, C.-S. Jiang, J. van de Lagemaat, K. Zhu, Dograin boundaries dominate non-radiative recombination in CH3NH3PbI3perovskite thin films? Phys. Chem. Chem. Phys. 19 (2017) 5043e5050.

[62] Y. Yang, M. Yang, David T. Moore, Y. Yan, Elisa M. Miller, K. Zhu, MatthewC. Beard, Top and bottom surfaces limit carrier lifetime in lead iodideperovskite films, Nat. Energy 2 (2017) 16207.

[63] N. Nickel, N. Johnson, W. Jackson, Hydrogen passivation of grain boundarydefects in polycrystalline silicon thin films, Appl. Phys. Lett. 62 (1993)3285e3287.

[64] M.L. Terry, A. Straub, D. Inns, D. Song, A.G. Aberle, Large open-circuit voltageimprovement by rapid thermal annealing of evaporated solid-phase-crystallized thin-film silicon solar cells on glass, Appl. Phys. Lett. 86 (2005)172108.

[65] C.-S. Jiang, R. Noufi, J. AbuShama, K. Ramanathan, H. Moutinho, J. Pankow,M. Al-Jassim, Local built-in potential on grain boundary of Cu(In, Ga)Se2 thinfilms, Appl. Phys. Lett. 84 (2004) 3477e3479.

[66] Q. Chen, H. Zhou, T.-B. Song, S. Luo, Z. Hong, H.-S. Duan, L. Dou, Y. Liu, Y. Yang,Controllable self-induced passivation of hybrid lead iodide perovskites towardhigh performance solar cells, Nano Lett. 14 (2014) 4158e4163.

[67] H.J. Snaith, A. Abate, J.M. Ball, G.E. Eperon, T. Leijtens, N.K. Noel, S.D. Stranks,J.T.-W. Wang, K. Wojciechowski, W. Zhang, Anomalous hysteresis in perov-skite solar cells, J. Phys. Chem. Lett. 5 (2014) 1511e1515.

[68] E. Unger, E. Hoke, C. Bailie, W. Nguyen, A. Bowring, T. Heumüller,M. Christoforo, M. McGehee, Hysteresis and transient behavior in


currentevoltage measurements of hybrid-perovskite absorber solar cells,Energy Environ. Sci. 7 (2014) 3690e3698.

[69] Z. Xiao, Y. Yuan, Y. Shao, Q. Wang, Q. Dong, C. Bi, P. Sharma, A. Gruverman,J. Huang, Giant switchable photovoltaic effect in organometal trihalideperovskite devices, Nat. Mater. 14 (2015) 193e198.

[70] W. Tress, N. Marinova, T. Moehl, S. Zakeeruddin, M.K. Nazeeruddin, M. Gr€atzel,Understanding the rate-dependent JeV hysteresis, slow time component, andaging in CH3NH3PbI3 perovskite solar cells: the role of a compensated electricfield, Energy Environ. Sci. 8 (2015) 995e1004.

[71] R. Gottesman, P. Lopez-Varo, L. Gouda, J.A. Jimenez-Tejada, J. Hu, S. Tirosh,A. Zaban, J. Bisquert, Dynamic phenomena at perovskite/electron-selectivecontact interface as interpreted from photovoltage decays, Chem 1 (2016)776e789.

[72] C. Eames, J.M. Frost, P.R. Barnes, B.C. O'regan, A. Walsh, M.S. Islam, Ionictransport in hybrid lead iodide perovskite solar cells, Nat. Commun. 6 (2015)7497.

[73] J.M. Azpiroz, E. Mosconi, J. Bisquert, F. De Angelis, Defect migration inmethylammonium lead iodide and its role in perovskite solar cell operation,Energy Environ. Sci. 8 (2015) 2118e2127.

[74] S. Meloni, T. Moehl, W. Tress, M. Franckevi�cius, M. Saliba, Y.H. Lee, P. Gao,M.K. Nazeeruddin, S.M. Zakeeruddin, U. Rothlisberger, Ionic polarization-induced currentevoltage hysteresis in CH3NH3PbX3 perovskite solar cells,Nat. Commun. 7 (2016) 10334.

[75] J. Haruyama, K. Sodeyama, L. Han, Y. Tateyama, First-principles study of iondiffusion in perovskite solar cell sensitizers, J. Am. Chem. Soc. 137 (2015)10048e10051.

[76] E. Mosconi, D. Meggiolaro, H.J. Snaith, S.D. Stranks, F. De Angelis, Light-induced annihilation of Frenkel defects in organo-lead halide perovskites,Energy Environ. Sci. 9 (2016) 3180e3187.

[77] J.S. Yun, J. Seidel, J. Kim, A.M. Soufiani, S. Huang, J. Lau, N.J. Jeon, S.I. Seok,M.A. Green, A. Ho-Baillie, Critical role of grain boundaries for ion migration informamidinium and methylammonium lead halide perovskite solar cells, Adv.Energy Mater. 6 (2016) 1600330.

[78] J. Xing, Q. Wang, Q. Dong, Y. Yuan, Y. Fang, J. Huang, Ultrafast ion migration inhybrid perovskite polycrystalline thin films under light and suppression insingle crystals, Phys. Chem. Chem. Phys. 18 (2016) 30484e30490.

[79] Q. Wang, B. Chen, Y. Liu, Y. Deng, Y. Bai, Q. Dong, J. Huang, Scaling behavior ofmoisture-induced grain degradation in polycrystalline hybrid perovskite thinfilms, Energy Environ. Sci. 10 (2017) 516e522.

[80] C.H. Chiang, C.G. Wu, Film grain-size related long-term stability of invertedperovskite solar cells, ChemSusChem 9 (2016) 2666e2672.

[81] K. Domanski, B. Roose, T. Matsui, M. Saliba, S.-H. Turren-Cruz, J.-P. Correa-Baena, C.R. Carmona, G. Richardson, J. Foster, F. De Angelis, Migration ofcations induces reversible performance losses over day/night cycling inperovskite solar cells, Energy Environ. Sci. 10 (2017) 604e613.

[82] D. Kiermasch, P. Rieder, K. Tvingstedt, A. Baumann, V. Dyakonov, Improvedcharge carrier lifetime in planar perovskite solar cells by bromine doping, Sci.Rep. 6 (2016) 39333.

[83] M. Yang, T. Zhang, P. Schulz, Z. Li, G. Li, D.H. Kim, N. Guo, J.J. Berry, K. Zhu,Y. Zhao, Facile fabrication of large-grain CH3NH3PbI3-xBrx films for high-efficiency solar cells via CH3NH3Br-selective Ostwald ripening, Nat. Com-mun. 7 (2016) 12305.

[84] N.D. Marco, H. Zhou, Q. Chen, P. Sun, Z. Liu, L. Meng, E.-P. Yao, Y. Liu,A. Schiffer, Y. Yang, Guanidinium: a route to enhanced carrier lifetime andopen-circuit voltage in hybrid perovskite solar cells, Nano Lett. 16 (2016)1009e1016.

[85] W. Ke, C. Xiao, C. Wang, B. Saparov, H.S. Duan, D. Zhao, Z. Xiao, P. Schulz,S.P. Harvey, W. Liao, Employing lead thiocyanate additive to reduce thehysteresis and boost the fill factor of planar perovskite solar cells, Adv. Mater.28 (2016) 5214e5221.

[86] N.K. Noel, A. Abate, S.D. Stranks, E.S. Parrott, V.M. Burlakov, A. Goriely,H.J. Snaith, Enhanced photoluminescence and solar cell performance via Lewisbase passivation of organiceinorganic lead halide perovskites, ACS Nano 8(2014) 9815e9821.

[87] D. Bi, C. Yi, J. Luo, J.-D. D�ecoppet, F. Zhang, S.M. Zakeeruddin, X. Li, A. Hagfeldt,M. Gr€atzel, Polymer-templated nucleation and crystal growth of perovskitefilms for solar cells with efficiency greater than 21%, Nat. Energy 1 (2016)16142.























































































































































































































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The role of grain boundaries in perovskite solar cellsyylab.seas.ucla.edu/papers/MTE-review.pdfThe role of grain boundaries in perovskite solar cells Jin-Wook Lee 1, Sang-Hoon Bae