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Cation exchange for thin film lead iodideperovskite interconversion†

Giles E. Eperon, Clara E. Beck and Henry J. Snaith*

We report a new technique for tuning the bandgap of hybrid

organic–inorganic halide perovskite materials. By dipping films

of methylammonium or formamidinium lead triiodide (MAPbI3 or

FAPbI3) in solutions of formamidinium or methylammonium iodide

(FAI or MAI) at room temperature, we are able to inter-convert

through cation exchange between perovskite materials, allowing

us to carefully tune the bandgap between 1.57 and 1.48 eV. We

observe uniform conversion through the entirety of the bulk film,

with no evidence for a ‘‘bi-layered’’ or graded structure. By applying

this technique to solar cell devices, we are able to enhance the

performance of the single cation devices. Furthermore, we demon-

strate that this technique allows us to form pure black phase FAPbI3

infiltrated into mesoporous scaffolds; this is normally infeasible

since the pores confine the FAPbI3 in a yellow non-perovskite phase

with a much wider bandgap, which is not of practical use in solar

cells or other optoelectronic devices. Additionally, this work provides

evidence for molecular cation mobility in the halide perovskites,

indicating that the cations play a role in ionic conduction as well as

the mobile anions.

Introduction

Hybrid organic–inorganic lead halide perovskite solar cellshave recently attracted a lot of attention due to their meteoricrise in power conversion efficiencies, rising from 3.8% to over20% in a matter of years.1–4 These materials show promise notonly as solar cells but also as easily and cheaply fabricated semi-conducting materials for LEDs, lasers and transistors.5–9

The most widely studied organic–inorganic perovskite material,with a vast array of processing methods becoming prevalent,is methylammonium lead triiodide (MAPbI3).10 This perovskitehas a bandgap of B1.57 eV and can be formed at low tempera-tures (B100 1C). However, a number of recent reports have

studied a new perovskite, formamidinium lead triiodide (FAPbI3).This perovskite has the same ABX3 structure as MAPbI3 but anarrower bandgap, achieved by replacing the methylammoniumcation with a slightly larger formamidinium cation.2,11,12 It has abandgap of B1.48 eV, closer to the single-junction solar celloptimum, but it requires temperatures of 4150 1C to form therequired perovskite phase, residing in a yellow non-perovskitephase at temperatures below this.13 In fact, the optimum com-position for a perovskite solar cell may be a mixture of FAPbI3

and MAPbI3. By changing the ratio of MA and FA cations present,it has been shown that the bandgap can be tuned between that ofMAPbI3 and FAPbI3, and that these mixed compositions havefavourable properties in terms of structural stability in the blackphase and performance.2,13,14

Cation exchange is a well-studied concept in the field ofnanocrystals, for example the exchanges of Cu, Cd, Pb, In and

Department of Physics, Clarendon Laboratory, University of Oxford,

Oxford OX1 3PU, UK. E-mail: h.snaith1@physics.ox.ac.uk

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5mh00170f

Received 18th August 2015,Accepted 2nd October 2015

DOI: 10.1039/c5mh00170f

www.rsc.li/materials-horizons

Conceptual insightsHalide perovskite materials, solution-processable semiconductors withan ABX3 crystal structure, have recently become the subject of extremeinterest for their use in photovoltaics and other optoelectronics. The best-performing perovskite material, formamidinium lead iodide (FAPbI3), ismore difficult to process and requires higher temperatures than themore-studied wider bandgap methylammonium lead iodide (MAPbI3).In this work it is shown that bulk A-site cation exchange can take place inthese films when immersed in a solution containing a suitablereplacement cation, allowing conversion from MAPbI3 to the FAPbI3

(and in reverse) easily and at room temperature. Conversion takes placevia crystalline mixed-cation intermediates, allowing fine tuning of thebandgap. Importantly, by converting from an already-formed perovskite,this process allows the formation of the formamidinium perovskite blackphase in a mesoporous scaffold for the first time, a structure which hasdisplayed good results for the methylammonium material but normallyconstrains the formamidinium material in a non-perovskite yellow phase.A second key insight is that despite the bulky size of the cation, thematerials allow ion diffusion easily enough through the whole film toconvert the film uniformly with no gradation. Furthermore, solar cellsfabricated with mixed-cation materials made in this way show superiorperformance to the neat materials.

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Zn in nanocrystals has been demonstrated and is well-understood,often resulting in materials with properties superior to the initialparent material.15,16 After the synthesis of nanocrystals, they areexposed to a solution of the cation to be substituted, and over timethe cation exchange occurs. It can be driven by a preferentialchemical reaction or simple diffusion with mass action ensuringthe desired conversion is achieved. It is most often used to retainthe initial structure of a material whilst changing the compositionto one which cannot easily be obtained straight from precursors.In bulk films, ionic exchange is less well studied. There has beensome study on the insertion of halide anions into layered oxideperovskites.17 Most recently, Pellet and co-authors have reportedthe observation of rapid anion exchange in thin perovskite films,by exposing MAPbX3 (X = Cl, Br, I) to MAX solution. They observethat the halides are easily transported within the perovskite lattice,and that the morphology is generally preserved. However, theredoes not exist any study on whether transformation can occur inthe same way between the larger organic cations.

First principles calculations suggest that both anions andmolecular cations should be mobile species with low activation

energy for migration (0.6 eV for I� and 0.8 eV for MA+), takingpart in ionic conduction in the material, which suggests thatcation exchange should be able to occur in a similar manner tothe anion exchange.18

Herein we study the substitution of organic cations withinthe APbI3 perovskite structure (A = MA or FA). We expose aninitial perovskite film of MAPbI3 or FAPbI3 to solutions contain-ing a different cation (MAI or FAI) and use X-ray diffraction,absorption spectra, and photoluminescence to monitor conver-sion of the perovskite structure between MAPbI3 and FAPbI3.We find that on a time-scale of minutes, even quite thick filmsof perovskite, relevant for solar cells, can undergo conversionbetween perovskites, with mixed cation perovskites being formedin between. This allows us to overcome a previous limitation ofthe formamidinium based perovskite – forming it on a meso-porous scaffold, which previously has resulted in the yellowphase being formed, even with high temperature annealing.Furthermore, this provides a route for forming black phaseFAPbI3 at lower temperatures than normally needed. We alsofind that the presence of the mesoporous scaffold, as was found

Fig. 1 MAPbI3 conversion to FAPbI3�MAPbI3 films dipped in FAI solution for varying times. (a) Absorbance spectra showing onset of absorption.(b) Normalised X-ray diffraction spectra, magnified around 141. (c) Plot of peak position against dipping time, fit with an exponential decay. (d) Normalisedphotoluminescence spectra. (e) Scanning electron micrographs of films dipped for different times. Scale is the same in all cases.

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for anion exchange,19 vastly accelerates the rate of perovskiteinter-conversion. We fabricate solar cells with the convertedperovskites, and show that in some cases we can enhance photo-voltaic properties of the perovskite films by a simple dip treatmentof the as-formed film.

Results and discussion

Initially we studied the conversion of MAPbI3 by FAI exposure,dissolved in propan-2-ol. We formed MAPbI3 films via the non-stoichiometric chloride precursor route, to give films of B400 nmthickness, which were subsequently dipped at room temperaturein a dry atmosphere in 10 mg ml�1 FAI in propan-2-ol. Afterdipping for a certain time, films were rinsed with propan-2-oland dried at 100 1C. As we show in Fig. 1a, we found that upondipping the perovskite film into the solution, within tens ofminutes the absorbance begins to red-shift towards the FAPbI3

absorption onset, with a gradual shift from the neat MAPbI3

absorption towards a reference sample of neat FAPbI3, with theconversion slowing after around 240 minutes. After a signifi-cantly longer time (840 min), the red shift ceases to progressfurther. This implies that we are converting MAPbI3 to FAPbI3

via a cation exchange.We also show measured X-ray diffraction spectra in Fig. 1b.

FAPbI3 (black phase) has been recently reported to have a cubicstructure, whereas MAPbI3 has a tetragonal structure, at roomtemperature, the difference being a slight rotation of the leadiodide octahedra.20 However, these crystal structures are funda-mentally very similar so would be likely to support intermixingand facile conversion between the two without much latticestress. We show the full spectra in the ESI,† (Fig. S2); in Fig. 1bwe show a magnification around the peak observed at B141,which corresponds to the (110) peak of MAPbI3 tetragonalphase and the (100) peak of FAPbI3 cubic black phase. Heretoo we observe a continuous shift between the crystal struc-tures. Importantly there is no time at which we do not observe astrongly crystalline material, which might be expected if wehave too much lattice strain. This observation implies that wecan easily substitute the cations throughout the lattice withoutdetriment to the crystal structure. Moreover, we do not observetwo distinct peaks in any case, which we would expect toobserve if the film was segregated into MAPbI3 and FAPbI3

domains. Hence, the conversion must take place throughoutthe whole film and homogenise on the time scale of the experi-ment; the cations must easily move through the lattice. We notethat even after 840 minutes, the peak is not completely shiftedto the position of neat FAPbI3. As such, it appears that conver-sion is not 100% complete even after a very long time. Thisindicates that there is a thermodynamically stable point in thesubstitution of MA with FA, beyond which the reaction will notprogress further. We also observe an emergence of a peakcorresponding to the FAPbI3 yellow phase (see spectra in ESI†)at B11.81, arising after 120 minutes. To estimate a rate ofreaction, we plotted the position of the XRD peak at B141 as afunction of time (Fig. 1c); by fitting with an exponential decay

(as expected for a chemical substitution reaction) we extract atime constant of 110 minutes.19

We also measured photoluminescence (PL) of the convertedfilms. All films exhibited strong photoluminescence. Here, wealso observe a gradual transition from MAPbI3 PL to FAPbI3 PLspectra, in good agreement with the XRD and absorbance spectra,again indicating that the transition between perovskites occurs viafully crystalline mixed cation phases. In order to investigate themorphology, we took scanning electron micrographs, which weshow in Fig. 1e, of representative films at different points duringthe conversion. Initially the MAPbI3 has a reasonably continuousmorphology with large crystallites and a few large pinholes. As theconversion progresses, we observe little change in the macro-scopic structure, but we do observe the formation of someneedle-like crystals on the surface. The appearance of signifi-cant needle-like structures is coincident with the rise of peakscorresponding to the yellow phase in the XRD (see ESI†), so asproposed for a similar observation by Binek et al., we assignthese to the yellow phase of FAPbI3.13 This is the preferredcrystal habit of the yellow FAPbI3, and yellow FAPbI3 is thepreferred phase at room temperature, so it may assume thisphase and shape after a long enough time in the presence of apartial solvent (IPA) at room temperature.21

We next studied the opposite conversion: FAPbI3 to MAPbI3.We show data from the equivalent experiments in Fig. 2. Filmswere spin-coated from 1 : 1 FAI : PbI2 solution with hydroiodic acidadded before spin-coating, to enhance solubility and enable theformation of a smooth film, crystallised at 170 C for 10 minutesand then dipped in 10 mg ml�1 MAI in propan-2-ol, rinsing anddrying as before after dipping for a certain time.11 The films arealso about 400 nm in thickness. We observe in this case a blue-shift of the absorbance onset towards the MAPbI3 perovskiteonset, again indicating a continuous tuning of the bandgap,and that we are able to convert the perovskite in the same wayvia cation exchange. Fig. 2b and c show the X-ray diffractionpeak positions near 141 as a function of dipping time. We observehere, that we are also not able to fully convert the material toMAPbI3, with conversion getting closer to the reference datathan was the case for the conversion to FAPbI3, but not totallyreaching it. Fitting the peak position with an exponential fitagain, we extract a time constant of 36 minutes.

Although the absorbance and XRD data show a continuousshift between FA and MA materials here, the photolumines-cence spectra (Fig. 2d) behave in an unexpected manner. Upondipping in MAI, instead of observing a blue-shift towards theMAPbI3 photoluminescence peak, the photoluminescenceintensity immediately reduces and we observe it to be blue-shifted even further than the MAPbI3 spectrum, to B750 nmmaximum. Longer dip time does not appear to change this.This is clearly in discrepancy to the absorbance and XRDspectra. Such blue-shifted photoluminescence, which is notablybroader in energy than the PL peaks of either neat compounds,could arise from a strained perovskite phase, or some otherstructural disorder which increases the energetic disorder,where the broadening of sites occurs at higher energy thanthe band gap.22–25

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In SEM images, as we show in Fig. 2e, here the morphologydoes significantly change in the conversion from FAPbI3 toMAPbI3. Starting as an apparently flat film of small crystallites,with a few pinholes, upon dipping, the film initially appears tobecome more uniform with pinholes closing, but we observemore crystal edges and the crystals appear to be re-orienting.On longer times, we observe a definite change in crystal orienta-tion, appearing to form individual cuboid-like crystallites, butthe re-orientation results in more pinholes being formed. Thisis likely due to the crystals being afforded some mobility by thepropan-2-ol environment, in which the organic component issoluble. Over a long time, they are re-orienting to the preferredcrystal shape and size of the MAPbI3. As discussed by Stoumposet al., FAPbI3 and MAPbI3 have different crystal phases (forblack phase FAPbI3) and hence different preferred crystalhabit.21 MAPbI3 tends to an elongated rhombic dodecahedronbut FAPbI3 a regular rhombic dodecahedron. This physicalrearrangement of crystal shape appears to be responsible forthe initial closing followed by eventual appearance of pinholesin the film. We propose that this may also be responsible for

the blue-shifted PL, but much further work is required tounderstand the blue-shift in detail.

Having shown that we are able to convert between FAPbI3

and MAPbI3 flat films via cation exchange, we then turned tosolving a unique problem of FAPbI3 deposition. If FAPbI3 isspin-coated straight onto a mesoporous scaffold (thicker thanB200 nm), we find that it does not form the black a-phase evenon heating at high temperatures, as it does in a thin solid film.It remains in the yellow d-phase, appearing to be confined inthat crystal structure by the scaffold (see ESI,† Fig. S1, fordetails). Mesostructured perovskite cells remain heavily inves-tigated and may prove to be an important architecture, whethermesoporous TiO2 or mesoporous alumina are employed. Wenote that a thin mesoporous layer, with a thick capping layer ofperovskite on top, is able to form the black phase; it is when theperovskite is significantly infiltrated into a scaffold that weobserve the inhibition of black phase formation.2,12 Thus, if wecan form MAPbI3 infiltrated into a mesoporous scaffold andthen convert it into black phase FAPbI3 via this interconversionapproach, we could have a route to forming infiltrated black

Fig. 2 FAPbI3 conversion to MAPbI3�FAPbI3 films dipped in FAI solution for varying times. (a) Absorbance spectra showing onset of absorption.(b) Normalised X-ray diffraction spectra, magnified around 141. (c) Plot of peak position against dipping time, fit with an exponential decay. (d) Normalisedphotoluminescence spectra. (e) Scanning electron micrographs of films dipped for different times. Scale is the same in all cases.

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FAPbI3 perovskite. Since FAPbI3 has a lower bandgap thanMAPbI3, the conversion of MAPbI3 to FAPbI3 in a mesoporousscaffold, where we retain the same crystal shape and size due tothe constraints of the pores, should exhibit higher short-circuitcurrent and hence performance. To test whether this is possible,we carried out the same dip-conversion on MAPbI3 film infiltratedwithin mesoporous alumina of B400 nm thickness.

We show absorbance of interconverted mesoporous-infiltratedMAPbI3 films in Fig. 3a, compared to a reference of planar FAPbI3.We observe that after much shorter times than for the planarfilms, the absorbance onset shifts towards the FAPbI3 onset.Within 30 minutes of dipping time, the onset is at the sameplace as the FAPbI3 onset, indicating that almost full conver-sion has occurred. However, the total absorbance in this regionof the spectrum drops beyond 10 minutes of dipping time,indicating that there is likely some degradation or phase con-version happening. A magnification of the XRD spectra around141 shows that we do achieve very nearly full conversion fromMAPbI3 to black phase FAPbI3, and we note that the peaknarrows as dip time progresses. This indicates that crystallitesize is increasing; by 60 minutes it is almost as narrow as theplanar FAPbI3 spectrum. This suggests that the perovskite is nolonger being confined by the mesoporous scaffold (which limitscrystallites to B50 nm). To further investigate what is occur-ring, we plot the XRD spectra on a larger 2y scale, showing apeak associated with the yellow d-phase at 11.81. We observethat after B30 minutes, a significant (compared to the blackphase peak) amount of yellow phase FAPbI3 becomes apparent.We thus propose that over time, in propan-2-ol, which is apartial solvent, the perovskite rearranges to form yellow phase

crystals on top of the mesoporous scaffold, rather than inside it.These crystals are no longer confined by the mesoporous scaffoldand hence form larger crystallites than those confined by thescaffold. We fitted an exponential decay to the B141 peak posi-tion over time, and find a good fit with a time constant (t)of 14.7 minutes. Compared to the conversion of planar films(t B 117 minutes), this is much faster.

We can thus make some conclusions about the dynamics ofthe interconversion reaction. Firstly, the presence of a meso-porous scaffold accelerates the reaction – comparing MAPbI3

conversion we observe a t of 14.7 rather than 117 minutes.Presumably the larger surface area of perovskite infiltrated intothe mesoporous scaffold facilitates a faster reaction. This is ingood keeping with the observations of Pellet et al., who observeda similar process with anion exchange.19 Secondly, we observethat for similar thickness planar films, the reaction from FAPbI3

to MAPbI3 is faster than vice versa, with t of 36 rather than117 minutes. Whilst we note that the differing morphologies of thetwo perovskite films are likely to have an effect, this could also bedue to the reaction favouring an energetically preferred configu-ration. In this case, this would imply that the MAPbI3 perovskite isenergetically more favourable than FAPbI3. As recently pointed outby Binek et al., the interaction of the MA cation with the leadiodide lattice is stronger than the FA cation, since it has a greaterdipole moment (2.3D compared to 0.21D).26 This means thatstronger hydrogen bonds to the iodide ions will be formed inMAPbI3 than FAPbI3, making it a more energetically favourablelattice. The conversion to FAPbI3 will still be driven by the vastexcess of FA cations in the solution, but more slowly than theequivalent MA reaction. In all cases, almost total conversion

Fig. 3 MAPbI3 conversion to FAPbI3 in a mesoporous scaffold. (a) Absorbance spectra showing onset of absorption. (b) Normalised X-ray diffractionspectra, magnified around 141. (c) Magnification of XRD spectra around a larger region, showing the yellow d-phase region. (d) Plot of peak positionagainst dipping time, fit with an exponential decay.

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is achieved, with no bi-layer perovskites being formed. As such,the mobility of cations throughout the perovskite lattice mustbe rapid and facile. This observation suggests that the assertionof Lee and co-workers, that by spin-coating MAI in propan-2-olon top of FAPbI3 they could form a layer of MAPbI3 as a discretelayer on top of FAPbI3, is likely to be incorrect.12 This techniqueis analogous to the dip-conversion and we never observe twodistinct crystal structures for MAPbI3 and FAPbI3.

Having demonstrated the conversion between different pero-vskites, we fabricated solar cells in order to investigate the impactof this conversion on devices. We note that there are several effectsat play in the conversion reactions; as well as the conversion fromone perovskite to another, we have the introduction of strain onthe lattice, the changing morphology of the planar films, and thedegradation into the yellow FAPbI3 phase at long dipping times.

We fabricated planar solar cells in the architecture FTO/compact TiO2/perovskite/spiro-OMeTAD/Au, and mesoporous

solar cells with the same architecture but with a layer of meso-porous alumina deposited prior to the perovskite deposition.We deposited the perovskites layers as for the fabrication offilms. We show current–voltage characteristics for championdevices, and average PCEs for the batch, at each dipping time inFig. 4. We note that these devices did display current–voltagehysteresis, and stabilised PCEs are shown in the ESI,† (Fig. S4–S6).Considering first the MAPbI3 to FAPbI3 interconversion (Fig. 4aand b), we observe that for the first 2 hours, performancemonotonically decreases slightly, mainly due to a reductionin Voc, and after that both Jsc and Voc are reduced significantlyleading to poorly-performing devices for the longest dippingtimes. Though there is some evidence for a small increase inJsc between 0 and 30 min, as we might expect for a narrowerbandgap material, the performance is dominated by thereduction in Voc so there is no net gain in efficiency. Gradualformation of the yellow phase FAPbI3 needle structures is likely

Fig. 4 Solar cell characteristics of interconverted perovskite solar cells. Current–voltage characteristics measured under AM1.5 illumination for(a) MAPbI3 conversion to FAPbI3 (planar), (c) FAPbI3 to MAPbI3 (planar) and (e) MAPbI3 to FAPbI3 on mesoporous alumina scaffolds. Average fast-scanned PCE determined from J–V curves is shown in (b), (d), and (f) respectively. Devices were held at 1.4 V for 5 s before measuring, and measuredfrom 1.4 V to 0 V at 0.38 V s�1. Average data is taken from at least 7 devices per point.

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to be the cause of the dramatic drop in performance at longconversion times.

Considering then the FAPbI3 to MAPbI3 interconversion(Fig. 4c and d), we observe that upon dipping for 10 minutesor more, the PCE increases (except for an apparently anomalous30 minute dipping result). This is due mainly to an increase inshort-circuit current. We note that this is not entirely expected; awider bandgap material (MAPbI3) should produce less current.However, looking more closely, we note that at 10 minutes thebandgap has not shifted much (Fig. 2a and b), but looking at theSEM images (Fig. 2e), the film becomes more continuous, withfewer pinholes. This will also have the effect of increasing short-circuit current, simply because there are fewer regions wherelight can pass straight through. We propose that this morpholo-gical effect, rather than any intrinsic electronic material property,is dominating the solar cell performance change. The perfor-mance begins to drop again at the later dipping times, and this isin keeping with the again reduced uniformity of the perovskitefilm. We note that the unexpectedly blue-shifted PL has noobvious detrimental effect on device performance here.

Finally, we consider the mesoporous devices, converted fromMAPbI3 to FAPbI3. Upon short dipping times, up to B5–10 minutes,we observe higher PCE. In particular, for this batch a 2 minute dipshows significantly improved PCE. We note that more experimentaldata would be necessary to conclude the optimum ‘perfect’ dippingtime, due to the variation between 2, 5 and 10 minutes points. Thisfull optimisation is beyond the scope of this study, however – herewe conclude that a short dipping time is able to give rise to higherPCE. Upon longer dipping times, performance drops off criticallyand by 30 minutes it is very poor. The improvement observed atshort dip times comes mainly from an enhanced short-circuitcurrent. Since at 10 minutes, we do observe a significant shift inmaterial bandgap to a narrower value (Fig. 3), it is likely thatthis is likely responsible for the short-circuit current improve-ment; morphology plays less of a role in mesostructured devices.At longer dip times however, the performance drops off again,corresponding to production of the yellow FAPbI3 phase whichwe observed in Fig. 3.

As such we can conclude that for solar cell devices, whilstmorphological and phase change effects play a crucial role indetermining device performance, we are able to employ interconver-sion via dipping in MAI or FAI in propan-2-ol to fabricate higherperforming devices. The optimum compositions appear to be some-where in between MAPbI3 and FAPbI3 in the cases where we observean improvement in device performance; this is in keeping withrecent reports on the highest efficiency perovskite devices.2

FAPbI3 normally requires heating at high temperatures(4150 1C) to form the black phase. While we note that we doform some yellow phase, we are able to form FAPbI3 blackphase by simply dipping a MAPbI3 film at room temperature ina solution of FAI, rinsing in propan-2-ol, and drying at 100 1C.This therefore reduces the fabrication temperature of a FAPbI3

film. While planar solar cells do not show improvements, due tomorphological considerations and the production of some yellowphase, this technique may be able to be further optimised to leadto improvements, and could be of use in the field of lasers and

LEDs. Moreover, while it is normally infeasible to form black phaseFAPbI3 infiltrated in a mesostructured scaffold architecture, wehave demonstrated that by inter-converting from infiltratedMAPbI3 we can form this. Whilst our solar cells do not performwell due to detrimental formation of some yellow phase, withfurther optimisation this could be of use in such devices. Mostimpressively, the material made by incomplete interconversion,a mixture of FAPbI3 and MAPbI3 (MAxFA1�xPbI3) shows improvedsolar cell performances compared to either neat phase. Thisimplies that by a simple and quick dipping treatment we canenhance solar cell properties, and this may lead the way towardseven higher efficiencies in the future.

An important aspect of this work is that it demonstratesstrong evidence for high mobility of molecular cations in thesematerials. As reported by several groups recently, these perovskitematerials appear to be good ionic conductors, but it was not knownwhich species were moving – it was assumed to be predominantlythe halide anions.18,27 However, it was also reported via firstprinciples calculations that the activation energy for migration ofthe MA+ species would be just 0.2 eV more than the I� species,indicating that cation mobility is also probably high.18 Here, weobserve comprehensive replacement of cations in the perovskitematerial, though at a slower rate than the anion exchange reportedby others.19 This suggests that the cations are indeed easily mobile,though less so than the anions, in good agreement with the higheractivation energy predicted. Ionic conduction in perovskites is thuslikely to involve both anions and molecular cations, with highermobility for the anions.

Conclusion

We have demonstrated that by dipping a film of FAPbI3 orMAPbI3 in MAI or FAI respectively, dissolved in propan-2-ol,organic cation exchange occurs, allowing us to convert almostfully between perovskites. We observe a transformation of the bulkmaterial rather than formation of a bi-layered or graded structure.During conversion, we form crystalline intermediate phases ofMAxFA1�xPbI3. This allows us to carefully tune the bandgap of thisperovskite between 1.57 eV and 1.48 eV, having application in solarcells, and light emitting applications. Besides simply tuningthe bandgap, we observe critical influences on changing themorphology and formation of yellow phase FAPbI3. Despite this,we are able to significantly enhance the performance of hybridperovskite solar cells made of pure phase materials by exposingthem to this simple dip-interconversion process. Importantly, thiswork provides good evidence for molecular cation mobility in thehalide perovskites, indicating that ionic conduction in thesematerials likely involves both anions and molecular cations.

ExperimentalMaterials

Unless otherwise stated, all materials were purchased fromSigma-Aldrich or Alfa Aesar and used as received. Spiro-OMeTADwas purchased from Borun Chemicals.

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Perovskite film fabrication

Formamidinium iodide (FAI) and methylammonium iodide (MAI)were synthesised in-house according to reported procedures.3,11

To form the FAPbI3 precursor solution, FAI and PbI2 weredissolved in anhydrous N,N-dimethylformamide (DMF) in a1 : 1 molar ratio, at 0.55 M of each reagent. Immediately priorto spin-coating, 38 ml of hydroiodic acid (57% w/w) was addedto 1 ml of the 0.55 M precursor solution to enable smooth andcontinuous film formation.11 To form films, the precursor wasthen spin-coated in a nitrogen-filled glovebox at 2000 rpm, thenannealed at 170 1C in air for 10 minutes.

To form the MAPbI3 precursor solution, methylammoniumiodide and PbCl2 were dissolved in DMF in a 3 : 1 molar ratio, at0.88 M PbCl2 and 2.64 M MAI. To form films, the precursor wasspin-coated at 2000 rpm in a nitrogen-filled glovebox, allowedto dry at room temperature in the glovebox for 30 minutes,then annealed at 90 1C for 150 minutes followed by 120 1C for15 minutes in the glovebox.

Films for optical studies were fabricated on plasma-cleanedmicroscope glass slides. For solar cells, films were fabricated onfluorine-doped tin oxide (FTO) coated glass (Pilkington, 7 O&�1).Initially FTO was removed from regions under the anode contactby etching the FTO with 2 M HCl and zinc powder. Substrateswere then cleaned sequentially in hallmanex detergent, acetone,propan-2-ol and oxygen plasma.

Inter-conversion

For the dip-interconversion, films were immersed in solutions ofMAI or FAI in propan-2-ol in a nitrogen-filled glovebox at a concen-tration of 10 mg ml�1 for varying times. After dipping, films wererinsed in propan-2-ol and then dried at 100 1C for 30 minutes.

Device fabrication

A B50 nm hole-blocking layer of compact TiO2 was firstdeposited by spin-coating a mildly acidic solution of titaniumisopropoxide in ethanol (350 ml in 5 ml ethanol with 0.013 MHCl) at 2000 rpm, and annealed at 500 1C for 30 minutes. For themesoporous cells, 400 nm mesoporous alumina was then deposi-ted by spin-coating at 2500 rpm a solution of 20 nm aluminananoparticles suspended in propan-2-ol, which was then dried at150 1C. Here, the perovskite was deposited as before but spin-coated and annealed in air at 100 1C for 120 minutes.

Following interconversion, the hole-transporting layer wasthen deposited via spin-coating a 0.0788 M solution in chloro-benzene of 2,20,7,70-tetrakis-(N,N-di-p-methoxyphenylamine)-9,90-spirobifluorene (spiro-OMeTAD), with additives of 0.0184lithium bis(trifluoromethanesulfonyl)imide (added in 0.61 Macetonitrile solution) and 0.0659 M 4-tert-butylpyridine. Spin-coating was carried out at 2000 rpm.

Gold electrodes were then thermally evaporated under vacuumof B10�6 Torr, at a rate of B0.1 nm s�1, to complete the devices.

Device characterisation

The current density–voltage ( J–V) curves were measured (2400Series SourceMeter, Keithley Instruments) under simulated AM

1.5 sunlight at 100 mW cm�2 irradiance generated by an AbetClass AAB sun 2000 simulator, with the intensity calibratedwith an NREL calibrated KG5 filtered Si reference cell. Themismatch factor was calculated to be 1.2% between 400 and1100 nm. The solar cells were masked with a metal aperture todefine the active area, which was 0.0919 cm�2, and measured ina light-tight sample holder to minimize any edge effects.

Optical measurements

Absorbance spectra were collected with a Varian Cary 300 UV-Visspectrophotometer with an internally coupled integrating sphere.

Materials characterization

A Hitachi S-4300 field emission scanning electron microscopewas used to acquire SEM images. Sample thicknesses weremeasured using a Veeco Dektak 150 surface profileometer.X-ray diffraction spectra were obtained using a Panalytical X’PertPro X-ray diffractometer.

Photoluminescence measurements

Steady-state PL measurements were acquired using a FluoTime300 (PicoQuant GmbH) system. Film samples were photoexcitedusing a 510 nm laser head (LDH-P-C-510). The PL was collectedusing a high resolution monochromator and hybrid photomulti-plier detector assembly (PMA Hybrid 40, PicoQuant GmbH).

Acknowledgements

This work was part funded by EPSRC and from the EuropeanUnion Seventh Framework Programme [FP7/2007-2013] undergrant agreement no. 604032 of the MESO project. GE was sup-ported by the EPSRC and Oxford Photovoltaics Ltd through aNanotechnology KTN CASE award. CB was supported by ZuhlkeEngineering AG. We thank Jin Zhang and Martina Congiu forcleanroom maintenance and materials synthesis.

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