Accelerated degradation due to weakened adhesion from Li-TFSI

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Accelerated degradation due to weakened adhesion from Li-TFSI additives in perovskite solar cells Inhwa Lee, Jae Hoon Yun, Hae Jung Son, and Taek-Soo Kim ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b14089 • Publication Date (Web): 01 Feb 2017 Downloaded from http://pubs.acs.org on February 3, 2017

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ACS Applied Materials & Interfaces

Accelerated degradation due to weakened adhesion from Li-TFSI additives in perovskite solar cells Inhwa Lee1, Jae Hoon Yun2, Hae Jung Son2, and Taek-Soo Kim1* 1

Department of Mechanical Engineering, KAIST, Daejeon, 305-701, Republic of Korea

2

Photoelectronic Hybrid Research Center, Korea Institute of Science and Technology (KIST), Seoul, Republic of Korea

KEYWORDS: Perovskite solar cell, moisture, adhesion, additive, decomposition mechanism

1

ACS Paragon Plus Environment


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ABSTRACT Reliable integration of organometallic halide perovskite in photovoltaic devices is critically limited by its low stability in humid environments. Furthermore, additives to increase the mobility in the hole transport material (HTM) have deliquescence and hygroscopic properties, which attract water molecules and result in accelerated degradation of the perovskite devices. In this study, a double cantilever beam (DCB) test is used to investigate the effects of additives in the HTM layer on the perovskite layer through neatly delaminating the interface between the perovskite and HTM layers. Using the DCB test, the bottom surface of the HTM layers is directly observed, and it is found that the additives are accumulated at the bottom along the thickness (i.e., through-plane direction) of the films. It is also found that the additives significantly decrease the adhesion at the interface between the perovskite and HTM layers by more than 60% through hardening the HTM films. Finally, the adhesion-based degradation mechanism of perovskite devices according to the existence of additives is proposed for humid environments.

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INTRODUCTION Photovoltaic devices based on organometallic halide perovskites (e.g. methylammonium lead halides CH3NH3PbX3, where X = halogen) have attracted significant attention as the fastest advancing solar cell technology with high power conversion efficiencies (PCEs)1,2. Unlike conventional solar cell technologies, perovskite solar cells can be fabricated using solution-based processes, which enable various advantages such as simple equipment, high fabrication speed, lower material consumption, and low cost3,4. However, perovskite solar cells have two critical drawbacks: the use of toxic lead which disrupts large-scale device fabrication5-7 and low longterm device stability due to decomposition by moisture8. Therefore, current research has focused on replacing Pb with other candidates such as Sn5,6 and Bi7, and also increasing water resistance. The previous work by Frost et al.9 demonstrated the potential decomposition pathway of the perovskite layer in humid environments reporting that volatile components were formed during the decomposition process. It is possible that the remaining volatile components and decomposed molecules cause thin film swelling or pressure between the thin films, and thus accelerate the delamination. However, the methods to improve the characteristics that are vulnerable to water have focused on employing protective layers10, finding appropriate hole transporting materials (HTMs)11,12, and changing the electrodes13-15 even though the stability of perovskite is highly related to adhesion according to the decomposition process. Moreover, Yun et al.16 highlighted the importance of strong adhesion between the perovskite and the HTM layers because a weak interface can initiate hazardous failure in the devices through causing delamination on a large scale. In addition, most HTMs for solution-process perovskite solar cells need additional ion additives for high hole mobility and device performance17-20. However, the hygroscopicity of the 3



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ion additives in HTMs attracts water molecules, which results in accelerated degradation. This is completely opposite to the intended function of the HTM layer as an encapsulation layer by protecting external water molecules. It is possible to enhance the efficiency of perovskite devices through using ion additives, but it will only be beneficial for short-term stability21,22. The work by Bi et al22 showed that the increased PCE of 12.3% with the addition of Li-TFSI was decreased to 3.6% when the device is exposed to a humidity of 60% for 150 h. Therefore, not only improving the efficiency of devices but also deepening the understanding of the decomposition of perovskite solar cells due to water molecules is needed for long-term stability based on the effect of ion additives on the mechanical characteristics. In particular, a specific approach in terms of the adhesion mechanism is required because the weak interface in multilayered perovskite solar cells can function as a degradation pathway for water molecules in the presence of water. In this work, the accelerated decomposition mechanism of perovskite solar cells is proposed through the weakened adhesion due to ion additives under humid environments. In order to precisely examine the effects of additives on the perovskite layer, a double cantilever beam (DCB) test was used to neatly delaminate the HTM layer from the crystalline perovskite layer. Using the DCB test, it could be determined that the additives significantly decrease the adhesion by

more

than

60%

for

the

2,2′,7,7′-tetrakis(N,N-di-p-methoxyphenyl-amine)-9,99-

spirobifluorene (spiro-OMeTAD) and poly(3-hexylthiophene-2,5-diyl) (P3HT) HTM materials, and the weakened adhesion at the interface causes large-scale delamination when exposed to water molecules. Furthermore, the bottom surface of the HTM layers was directly observed using x-ray photoelectron spectroscopy (XPS) analyses and it was found that the additives accumulated at the bottom along the thickness (i.e., through-plane direction) of the films. Finally, 4


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the correlations between the additives, adhesion, and delamination were investigated.

EXPERIMENTAL SECTION Materials. A titanium diisopropoxide bis(acetylacetonate) solution (75 wt% in isopropanol), n-butanol (99.8%), 2-propanol (anhydrous, 99.5%), N,N-dimethylformamide (anhydrous, 99.8%), lead(II) iodide (99%), 4-tert-butylpyrididne (tBP, 96%), chlorobenzene (anhydrous, 99.8%), and bis(trifluoromethane) sulfonamide lithium salt (Li-TFSI, 99.95%) were purchased from

Sigma-Aldrich.

2,2′,7,7′-tetrakis(N,N-di-p-methoxyphenyl-amine)-9,99-spirobifluorene

(spiro-OMeTAD, Merck), poly(3-hexylthiophene-2,5-diyl) (P3HT, Rieke Metals, Inc.), TiO2 paste (Dyesol 18NRT, Dyesol), methylammonium iodide (Dyesol), and gold (Au, 99.99%, Itasco) were used to fabricate the perovskite films. Sample Fabrication. Spiro-OMeTAD and P3HT were used as the HTM layers, which represent small molecules and polymers, and their chemical structures are presented in Figures 1a and b, respectively. Also, Figures 1c and d present the bis(trifluoromethane)sulfonamide lithium salt (Li-TFSI) and 4-tert-butylpyrididne (tBP) structures that were mixed with the HTM layers in order to increase the hole mobility. In order to examine the effect of the ion additives on the perovskite layer, multilayer structures were fabricated as described in Figure 1e. The specimens were varied with the HTM species and the existence of additives.

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Figure 1. Chemical structures of (a) spiro-OMeTAD, (b) P3HT, (c) Li-TFSI and (d) tBP; and (e) cross-sectional SEM image of a perovskite solar cell.

Fluorine-doped tin-oxide (FTO) substrates (TEC-8, Pilkington) were cleaned using ethanol, isopropanol, and acetone, sequentially. After the UV-O3 treatment for 20 min, a compact TiO2 layer was spin-coated onto the FTO substrate at 500 rpm for 5 s, 1000 rpm for 5 s, and 2000 rpm for 40 s successively from a titanium diisopropoxide bis(acetylacetonate) solution (75 wt.% in isopropanol) diluted with n-butanol (1:11, v/v). Then, the film was dried on a hotplate at 120 °C for 10 min, followed by annealing at 500 °C in a furnace for 30 min and cooling to room temperature. Then, the mesoporous TiO2 layers prepared from a commercially available TiO2 paste that was diluted in ethanol (2:7, w/w) were formed via spin-coating at 3500 rpm for 40 s and dried at 120 °C for 10 min in order to evaporate any residual solvents. Next, the films were transferred to the furnace and annealed at 500 °C for 1 h. The perovskite layer was then obtained through a sequential deposition method. Before spin-coating, a PbI2 solution (1.0 M in DMF) dissolved at 80 °C was maintained warm at 80 °C. Then, the PbI2 layer was deposited via spincoating at 6500 rpm for 1 min after a 30 s delay before spinning to infiltrate the solution into the mesoporous scaffold. The perovskite layer was formed after spin-coating a CH3NH3I solution (10 mg/ml in 2-propanol) at 1000 rpm for 10 s, followed by thermal annealing at 80 °C for 20 6


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min. The HTM solutions for spiro-OMeTAD and P3HT were prepared in chlorobenzene of 56 mg/ml for spiro-OMeTAD and 15 mg/ml for P3HT with additive-free HTMs; then, Li-TFSI (HTM/Li-TFSI ratio, 10:1, w/w) and t-BP (HTM/tBP ratio, 10:1, w/w) were added for the HTMs with additives. The hole transporting layer was formed via spin-coating at 2500 rpm for 20 s (spiro-OMeTAD) and 3000 rpm for 30 s (P3HT). Finally, a 100 nm layer of Au was thermally evaporated to cover the entire films. Double Cantilever Beam (DCB) testing. The DCB specimen was fabricated through cutting the 25.4 × 25.4 mm glass into a rectangle of 8 × 25.4 mm after the Au deposition process. Then, a dummy glass substrate of 8 × 25.4 mm was attached to the previously cut sample using epoxy to make a sandwiched structure as depicted in Figure 3a. A high precision micromechanical test system (Delaminator Adhesion Test System, DTS Company, USA) was used to measure the fracture energy. The sandwiched thin films between the glass substrates were under loading and unloading at a constant displacement rate of 0.5 µm s-1; the applied load versus displacement were recorded continuously. From the load versus displacement curve, the crack length (a) and the fracture energy (Gc) can be calculated using the following equations:

a = =

/

− 0.64ℎ,

1 + 0.64 ,

(1)

(2)

where C is the elastic compliance, E’ is the plane-strain modulus of the beam, B is the specimen width, h is the half height for the substrate, P is the applied load, and Pc is the critical load where the load of the load versus displacement curve begins to decrease. 7



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Surface characterization techniques. The surface morphologies of the delaminated HTM layers were investigated using an atomic force microscope (AFM, Park Systems XE-100, Korea) in the non-contact mode. Furthermore, nanoindentation (Nano Indentation System XP, MTS, USA) was used to measure the hardness of the films according to the ion additives for further mechanical analyses of the nanofilms. The measurement was conducted 9 times for each specimen and the hardness data were obtained between depth of 200 and 300 nm to avoid substrate effect which can affect the indentation results considering the tip size diameter of 150 nm.

RESULTS AND DISCUSSION Degradation in humid environments. In order to investigate how additives affect the degradation of perovskite solar cells, spiro-OMeTAD and P3HT with and without additives were exposed to humid environments of 85 % humidity at 25 °C. The optical images in Figure 2a depict that the additives accelerate the degradation of the perovskite/spiro-OMeTAD structure resulting in enlarged holes and highly color changed surfaces in the spiro-OMeTAD layers. P3HT also exhibited similar results when additives were added, as depicted in Figure S1. First, regardless of the addition of additives, the HTM layers had pinholes on the surfaces before exposure to the water molecules as indicated by the red circles in Figure 2a. However, for the HTM with additives, more pinholes were observed. The differences between the with and without additive conditions became clear when the pristine perovskite/spiro-OMeTAD structure was exposed to humid environments for 24 h. For that without additives, the degradation occurred around the pinholes of the HTM layer. However, for that with additives, the perovskite layer, which was located under the HTM layer, was exposed to water molecules after 24 h 8


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exposure. Furthermore, the film that mixed additives in the HTM had a larger size compared with that without additives. From the changes, it is noticeable that the additives significantly accelerated the degradation of the pristine perovskite/HTM structure and the phenomenon became severe with increases in the exposure time.

Figure 2. (a) Optical images of the spiro-OMeTAD specimens with and without additives obtained for the pristine sample, after 24 h exposure, and after 48 h exposure under humid environments of 85% humidity at 25 °C. Cross-sectional SEM images of 24 h exposed (b), (c) spiro-OMeTAD and (d), (e) P3HT specimens of with and without additives.

Both Li-TFSI and tBP additives degrade the long term stability of the perovskite devices. In case of Li-TFSI, the additive attracts water molecules and it results in highly degraded films due to the hygroscopicity of the ion additives while tBP additive makes perovskite layer corrosive from the reaction between tBP and PbI2.23 However, because our degradation mechanism is focused on the effects by humid environments, direct attractions of water molecules by Li-TFSI 9



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additives to perovskite layer are considered as a main degradation mechanism. In particular, this phenomenon caused by the Li-TFSI additive is harsh for perovskite solar cells because the perovskite layer reacts with water molecules. Previous research has categorized the degradation mechanisms into two processes of reversible24 and irreversible9 processes depending on the amount of water25. According to the reversible process24, the conversion from MAPbI3·H2O to MAPbI3 is possible when it is hydrated and dehydrated. However, in the irreversible process9, the water molecules cause decomposition of the MAPI into aqueous HI, solid PbI2, and CH3NH2 either released as gas or dissolved in the water. The key difference separating the reversible and irreversible degradation processes is the presence of condensed water. Considering the phenomenon, the degradation condition in this study is an irreversible process due to high humidity conditions of 85% at RT. Therefore, volatile organics and gases can be formed during the decomposition process using water molecules that might induce pressure at the interface between the perovskite and HTM layers. Furthermore, the adhesion can be significantly reduced during the decomposition process because the solid components of the perovskite that were in physical contact with the HTM layer change to gas. Therefore, it is expected that the decomposition process can cause large areas of delamination at the interface. However, the large-scale delamination was only observed in the condition with the additives in the spiroOMeTAD in the cross-sectional SEM images in Figure 2c (SEM image of low magnification is also included in Figure S2). Even when P3HT with additives was exposed to identical humid conditions, the large area delamination did not occur and only local regions were delaminated (Figure 2e). It is believed that the large-scale delamination is related to the adhesion at the interface according to the existence of additives and the HTM properties; thus, the interfacial fracture energy was further measured using a DCB test. 10


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Adhesion at the interface between the perovskite and HTM. In order to investigate the effect of the ion additives on the fracture energy at the interface between the perovskite and HTM layer, a DCB test was conducted and it systematically measured the fracture energies. DCB tests have been used for various thin films such as organic solar cells, nanoparticle thin films26,27, graphene28, and conductive polymer29 layers because the method is more quantitative compared with other techniques (e.g. the die shear test and the peel test). The measured critical fracture energy (Gc in J/m2) is defined as the macroscopic work of the fracture per unit area to separate two layers. For the DCB test, a sandwiched specimen structure was fabricated as depicted in Figure 3a with various HTM layer conditions of with and without ion additives. For the Au top electrode, there was no delamination problem. The reason why Au delamination did not occur is believed to be associated with the high work of adhesion at the interface between polymer and Au layers, and also with the crack initiation from the micro defects at the interface between perovskite and HTM layers. The work by Kendall et al30 reported that the ultrasmooth surfaces of mica, gold and polymers cannot be delaminated when they are closely in a contact due to the van der Waals forces. Because surfaces of HTM layers are expected to be ultrasmooth due to spin-coating, strong attraction between Au and HTM layers are possible. Also, the fracture paths can be initiated to the interface between perovskite and HTM layers due to the micro defects such as voids and cracks at the interface. The tests for measuring Gc were conducted at 25 °C and 40 % relative humidity.

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Figure 3. (a) Schematic of the DCB specimen for measuring the debonding energy. (b) Measured fracture energies of the spiro-OMeTAD and P3HT depending on the additive addition before exposure to humid environments. (c) and (d) SEM and AFM images of spiro-OMeTAD, respectively, and (e) and (f) SEM and AFM images of P3HT, respectively.

Figure 3b depicts the measured fracture energies of the spiro-OMeTAD and P3HT with and without ion additives, respectively. From the results, it is clear that the pristine P3HT exhibited 214.9% larger fracture energy of 5.07 ± 0.30 J/m2 than that of the pristine spiro-OMeTAD of 1.61 ± 0.30 J/m2. Figures 3c to 3f explain the reason for the P3HT layer exhibiting a significantly higher fracture energy. Figures 3c and 3d and Figures 3e and 3f present the SEM and AFM images for the delaminated HTM surfaces on the spiro-OMeTAD and P3HT without additives after the DCB tests, respectively. For both surfaces, pinholes were observed that resulted from the sharp and pointed regions of the perovskite surface. However, the surface morphologies varied highly depending on the species of HTM, because spiro-OMeTAD and P3HT have different characteristics with small molecules and polymer materials31, respectively. In general, 12


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polymer materials have highly entangled chains that result in large energy dissipations when cracks propagate due to the ductile characteristics and plastic deformation. In contrast, the interactions between the small molecules are very weak, so cracks are easily propagated. As a result, P3HT has a significantly rougher delaminated surface with an RMS of 40.49 nm compared with that of spiro-OMeTAD at 7.47 nm. Also, SEM images and XPS spectra are included for both sides of fractured spiro-OMeTAD and P3HT samples to support the fracture paths in Figure S3 and S4, respectively. In SEM images of Figure S3, the surfaces of perovskite side (Figure S3(a)(c)) has high resemblance with that of the perovskite. However, the surfaces are slightly covered by HTM because it is difficult to neatly delaminate the HTM from the perovskite layer that has a sharp and pointed surface. Therefore, although the fracture path follows the interface between perovskite and HTM layers, cohesive fracture occur at the interface due to the deeply penetrated HTM. Using XPS analysis (Figure S4), the cohesive fracture was also verified by detecting the key elements (Carbon, Oxygen and Nitrogen for Spiro-OMeTAD and Carbon, Sulfur for P3HT) of HTM for both sides. In addition to the difference of the fracture energies between the spiro-OMeTAD and P3HT, the decrease in the fracture energies is remarkable for the specimens with the ion additives added. Through adding the additives, the fracture energies of the spiro-OMeTAD and P3HT decreased to 0.65 ± 0.21 J/m2 and 1.76 ± 0.4 J/m2, respectively, from 1.61 ± 0.3 J/m2 and 5.07 ± 0.17 J/m2, respectively. Even though spiro-OMeTAD had weak adhesion due to the small molecule characteristics, the fracture energy was significantly decreased to under 1 J/m2. Here, the conditions were the same for the weakest fracture energy (Figure 3b) and the large scale delamination caused by the degradation under a humid environment (Figure 2b). For a value under 1 J/m2, it was reported that the value corresponds to a weak van der walls interaction 13



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without chemical bonding28. Therefore, it is expected that a fracture energy under 1 J/m2 is too weak to endure the pressure formed during the decomposition process and the transformation from a solid state to a gas state by the water molecules; thus, delamination occurs at the interface on a large scale. However, for P3HT, it is considered that even if additives significantly decrease the interfacial fracture energy, the adhesion is sufficient to prevent large-scale delamination through disrupting the additional crack propagation despite the occurrence of local region delamination. Analysis of the decreased fracture energies with ion additives. From Figure 3b, it is inferred that the ion additives affect the fracture energy through changing the mechanical characteristics of the HTM layers. In order to verify the effects of the additives on the adhesion, DCB specimens were fabricated without a perovskite layer in order to exclude the decrease in adhesion due to degradation by water molecules. Figure 4a depicts the measured fracture energy for the interface between the ITO and HTM with and without additives. Compared with the values of the fracture energies in Figure 3b, those of the fracture energies in Figure 4a are generally decreased because the perovskite layer, which causes the plastic deformation of HTM, was removed. Unlike the flat surface of ITO, the perovskite layer has a rough and pointed surface, which allows the HTM layers to penetrate deeper into the perovskite surface. Therefore, when crack propagates, more debonding energies are needed due to the mechanical interlocking and crack deflection. In particular, P3HT which has entangled chain structure deforms severely inducing energy dissipation from the mechanical interlocking and crack deflection during the delamination process, resulting in increased fracture energies. However, the decreasing trend of the fracture energy was identical when the additives were used. In order to analyze the reason for the decreased fracture energies due to the additives, nanoindentations were conducted on the 14


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fractured surfaces after the DCB test. The hardness results in Figure 4b depict that the hardness increased by 101.8 % when the ion additives were added. From the data, it is expected that the fracture energies decreased on the hard interfaces due to the ion additives. Through making the interface hard using the ion additives, plastic deformation was not possible during the delamination process. Therefore, the range of the decrease in the fracture energy was larger for P3HT because the material deformed more due to its polymer characteristic than the spiroOMeTAD with small molecules.

Figure 4. (a) Measured fracture energies of spiro-OMeTAD and P3HT according to the additive addition without a perovskite layer; (b) hardness results of the spiro-OMeTAD layer according to the existence of additives measured via nanoindentation; and (c) XPS analysis results on the top and bottom surfaces of the spiro-OMeTAD layer. One advantage of the DCB test is that it is possible to analyze the bottom surface of the HTM afterwards. Yu et al. also used the DCB test to reveal the interface between the electrode and electrolyte in fuel cells32. Figure 4c presents the XPS data for the top and bottom surfaces of the spiro-OMeTAD layer with the fluorine intensity, which is a key element of Li-TFSI additives. The bottom surface has a higher fluorine concentration of ~2.18 at.% compared with the top surface of ~0.49 at.%. From the data, it is assumed that the additives accumulated during the solution process. As stated above, the additives attract water molecules due to their hygroscopic 15



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properties and this results in the HTM layers being hard. Therefore, the accumulated additives at the bottom further deteriorate the perovskite layer, which was located immediately under the HTM layer, due to the low adhesion by the hard surfaces and highly attracted water molecules.

Degradation mechanism of perovskite layer under humid environments. The accumulated additives accelerate the degradation of perovskite solar cells through decreasing the fracture energies and attracting more water molecules to the interface between the HTM and perovskite layers. The mechanism for the accelerated degradation due to the additives was visualized without and with additives, as depicted in Figure 5. The schematic mechanism is based on the experimental results obtained at humid environments of 85 % humidity at 25 °C. Because higher temperature can accelerate the degradation process by humid environments, rapid formation of PbI2 and the loss of the organic component are expected.33,34 In the pristine HTM without additives (Figure 5a), pinholes exist on the HTM surface. Then, when the films are exposed to humid environments, the water molecules are concentrated at the pinholes (Figure 5b) through swelling the film along the pinholes. Even though the exposure time was longer, Figure 5c illustrates that the degradation focused on the HTM layers, not entire films, in perovskite solar cells.

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Figure 5. Degradation mechanism of perovskite solar cells under humid environments (a-c) without additives and (d-f) with additives.

However, when the Li-TFSI additives are added to the HTM, the additives accumulated at the bottom along the thickness of the HTM as depicted in Figure 5d. The dispersed additives in the HTM attract water molecules through the pinholes; thus, the entire HTM layer has low stability. In this case, the water molecules attack the interface between the HTM and perovskite due to the larger amount of additives at the bottom of the HTM (Figure 5e). Therefore, the perovskite layer becomes directly exposed to the humid environment and the reaction between the perovskite and water molecules occurs. From the reaction, methyammonia and volatile organic gases are formed with decomposed PbI2, as described below.

CH3NH3PbI3(s) !" CH3NH3I(aq) + PbI2 (s)

(3)

CH3NH3I(aq) ↔ CH3NH2(gas and aq) + HI (aq)

(4)

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The gases create pressure at the interface until the pressure is stabilized. At the same time, the solid components transform to gas states and thus delamination unavoidably occurs. However, if the adhesion is sufficiently high to disrupt additional crack propagation from the initial local region delamination, large-scale delamination can be prevented as described in Figure 2e. After the reactions, the gases begin to evacuate following the pinholes35,36. Therefore, the pinholes become larger according to the exposure time to the humid environment (Figure 5f). The work by Hawash et al.36 demonstrated that a higher concentration of Li-TFSI on the top surface in spiro-OMeTAD was observed after air exposure, which is unlike the freshly prepared film exhibiting a very low concentration of Li-TFSI. They indicated the pinholes can be migration paths that accelerate degradation. Furthermore, if the interface has a lower fracture energy (Gc) than the degradation (Gd) induced by the volatile organics and decomposed components during the decomposition process, the pressure opens the interface on a large scale; this highlights the importance of strong adhesion to prevent this. Measuring the fracture energy at the interface after exposure to the humid environment would quantify the relationship between adhesion and delamination. However, the humid environment makes the surfaces of HTM rough as shown in Figure 2a and thus it was impossible to deposit uniform top electrode on the specimens. However, recent work by Rolston et al37 reported the decreasing tendency in adhesion between doped P3HT and perovskite layers when exposed to 55% R.H. and 25 °C. From this work, it is expected that very low fracture energies would be measured from the delamination by volatile organics. From these results, in order to achieve reliable integration of the perovskite layer in photovoltaic devices, strong adhesion is required in order to maintain the efficiency and reliability from the trapped component. Furthermore, devices should be designed to reduce pinholes and to not use hygroscopic additives in the HTM. Therefore, recent works, such as the 18


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development of dopant free HTM, hydrophobic HTM and employing buffer layer between perovskite and HTM layers, have focused on avoiding the degradation caused by additives.

CONCLUSIONS The adhesion at the interface between the perovskite and HTM layers and the stability under humid environments were significantly affected by the existence of ion additives. The nanoindentation data demonstrated that the additives decreased the adhesion by causing the film to harden, and the XPS data confirmed that the hardness was more severe at the bottom of the HTM layer because the additives were accumulated along the thickness. Furthermore, the accumulated additives accelerated the attraction of water molecules to the perovskite, which results in highly degraded films with decomposed components. It is expected that the proposed degradation mechanism by the additives based on the adhesion will provide guidelines for the design of perovskite solar cells that satisfies both mechanical and electrical characteristics.

ASSOCIATED CONTENT Supporting Information Surface characterization of the P3HT specimens by using optical microscope

AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected] Notes The authors declare no competing financial interest. 19



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ACKNOWLEDGEMENTS This work was supported by the Wearable Platform Materials Technology Center (WMC) (2016R1A5A1009926), by the Basic Science Research Program (2015R1A1A1A05001115), and by the Global Frontier R&D Program in the Center for Multiscale Energy Systems (2014 M3A6A7074784) funded by the National Research Foundation under the Ministry of Science, ICT & Future Planning, Korea.

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Accelerated degradation due to weakened adhesion from Li-TFSI

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