J Llacer

D Moerman

O Douhéret

D Théron

P Leclère

" Addressing

Keywords: AFM Atomic force microscopy AM1, 5G Air mass 1, 5 global C-AFM Conductive atomic force microscopy

Global warming is one of the main concerns in our society nowadays, year after year the impact and consequences are becoming more visible. The Paris agreement set a target to limit the CO2 emission, which is mainly caused from the increasing demand for energy based on fossil fuels. Since then, the scientific community has increased their efforts on looking for clean energy sources such as renewable energies. In this regard, solar energy is meant to be one of the main energy sources that could replace fossil fuels. Therefore, photovoltaic technologies have evolved tremendously and, organic-inorganic hybrid halide perovskite materials have been one of the technologies with the fastest growth in solar cell performance. Perovskite materials possess unique properties such as flexibility, low-cost and ease to manufacture. Nonetheless, there are still some issues regarding their stability against atmospheric conditions that need to be understood. This dissertation focuses on the characterization of the electrical properties at the nanoscale of perovskitebased thin films by means of scanning probe microscopies (Conducting AFM, Kelvin Probe Force Microscopy, and Scanning Microwave Microscopy).

In this growing field of research, many perovskite structures, deposition methods, and synthesis routes have been developed and introduced in perovskite-based solar cells. In the first part of this dissertation, the context of perovskite materials is detailed and the methodology used through the thesis is also described. Then, we study and compare the electronic properties at the surface of perovskite materials synthesized following two different routes. Moreover, it is known that device engineering can increase both, the performance, and the stability of perovskite solar cells.

In a second part if the thesis, we show that for a given perovskite structure, the stability upon exposure to controlled small amounts of water can be significantly improved through the synthesis optimization. Finally, we provide a series of conclusions and perspectives that could help to further understand the perovskite behaviour at the local scale and to improve the cell performances.

Prologue

This thesis was carried out in a cotutelle between the University of Mons (Belgium) and the University of Lille (France). At the University of Mons, I was part of the Laboratory for Chemistry of Novel Materials (CMN). CMN is a research group focused mainly on modelling, synthesis and characterization of organic materials used for energy applications. The CMN is also founding the Materia Nova Research Center, located in Mons (Belgium). Materia Nova conducts a number of projects directed towards the industry and it allows the CMN department to extend their research activities. I had the opportunity to use the facilities from Materia Nova to synthesize and characterize the perovskite materials, which require an inert environment (usually a N2-filled glovebox) to work with. On the other side, in France I was working at the Institut d'Électronique, de Microélectronique et de Nanotechnologie (IEMN), located in Villeneuve d'Ascq (France). The IEMN is a research institute working in close collaboration with the University of Lille and the CNRS. Their research goes from nanoscience applications to instrumentation. At the IEMN, I had the chance to use their scanning microwave microscope and to test the perovskite samples previously made at Materia Nova.

First, I would like to thank Didier Théron and Philippe Leclère for giving me the opportunity to start this PhD. I offer my sincerest gratitude to Didier, for the sheer amount of knowledge you have imparted upon me. I hope I retained at least half of the things you have taught me. A substantial amount of thanks to Roberto Lazzaroni and Philippe, for all of the help and guidance you have provided me in order to accomplish the best results.

To a couple of guys I very much consider as my co-supervisors, Olivier Douhéret and David Moerman. Olivier, I thank you for the profound insights and conversations we had regarding my studies/results along this ever so long journey. I admire the energy you transmit to people on an everyday basis and I thank you for your patience that I have yet to see exhausted. David, to say that everything I know now about AFM is because of you, thanks for your guidance through my experimental work, you provided me with the tools and support I needed to crystallize this thesis. Also, I extend my gratitude to the all the Materia Nova members as well as to Petr Polovodov, who has been very supportive and helped me with lots of SMM measurements, thank you so much.

To Juan A. Anta and to Lidia Contreras Bernal for giving me the opportunity to work with you at the University Pablo de Olavide. Thank you both for your warm welcome in Sevilla and for teaching me about impedance spectroscopy.

To the friendly crew at the CMN department, while it does no justice to single people out, I will at least mention the people I shared priceless moments and/or trips. In no particular order: Pauline, Nadège, Alex, Séb, Max, Laura, Sinan and Fab, thank you all. I would like also to acknowledge some friendships I've made during my stay in Mons: Wouter, Luc and Mule. We've had some unforgettable moments, I will always treasure our time spent discussing about everything, playing squash, cycling and enjoying music. You were a strong support to me, I would never thank you enough.

To Paula C., doing a PhD in a country far from my friends/family is sometimes tough, but things get always easier with people like you, I am in awe of the support you have given me. A special thank you to my friend Joan, for a call when I needed it most. To Paula G., sharing with you the last 2 years of my PhD was like finding an oasis and will always be with me.

To my family, it would never be enough whatever I tell you here. Distance may separate us, but I always felt as if you were right next to me. I reckon with time you will forget what this paragraph meant so, there it goes: Tot el que soc és gràcies a vosaltres.

Résumé

Le réchauffement climatique constitue aujourd'hui une des principales préoccupations de notre société. Année après année, son impact et ses conséquences sont de plus en plus visibles. L'accord de Paris a fixé pour objectif de limiter les émissions de CO2, principalement issues de la demande croissante d'énergie à base de combustibles fossiles.

Depuis, la communauté scientifique a multiplié ses efforts dans la recherche de sources alternatives d'énergie propres telles que les énergies renouvelables. À cet égard, l'énergie solaire est censée être une des principales sources d'énergie susceptibles de remplacer les combustibles fossiles. Par conséquent, les technologies photovoltaïques ont énormément évolué et les matériaux, dont notamment ceux à base de pérovskites hybrides organiquesinorganiques halogénées, ont été parmi les technologies ayant connu la croissance la plus rapide des performances des cellules solaires. En effet, les matériaux de type pérovskite possèdent des propriétés uniques telles que la flexibilité, le faible coût et la facilité de fabrication. Néanmoins, il reste encore quelques problèmes importants concernant leur stabilité face aux conditions atmosphériques. Cette thèse porte sur la caractérisation des propriétés électroniques, à l'échelle nanométrique, des couches minces à base de pérovskite au moyen de microscopies à sonde locale (Conducting AFM, Kelvin Probe Force Microscopy, and Scanning Microwave Microscopy).

Dans ce domaine de recherche en pleine croissance, de nombreuses structures de pérovskite, des méthodes de dépôt et des voies de synthèse ont été développées et introduites dans des cellules solaires à base de pérovskite. La première partie de cette thèse consiste à détailler le contexte des matériaux pérovskites et à décrire la méthodologie utilisée dans le cadre de cette thèse. Ensuite, nous étudions et comparons les propriétés électroniques à la surface de matériaux pérovskites synthétisés suivant deux voies différentes. Il est également connu que l'ingénierie des dispositifs augmentait les performances et la stabilité des cellules solaires en pérovskite. Dans la seconde partie de cette thèse, nous montrons que, pour une structure de pérovskite donnée, la stabilité visà-vis de l'exposition contrôlée à des petites quantités d'eau peut être considérablement améliorée par l'optimisation de la synthèse. Pour terminer, nous proposons une série de conclusions et de perspectives qui pourraient aider à mieux comprendre le comportement de la pérovskite à l'échelle locale et à améliorer les performances des cellules solaires.

Chapter 1 Introduction

This chapter serves as an introduction to the evolution of photovoltaic technologies and specifically to hybrid perovskite materials used as solar cells. An overview of the perovskite properties, their performances in solar cells and the main concerns in the community is also included. Finally, the outline is provided along with the objectives of this thesis.

Photovoltaics

Energy is at the centre of the human civilization. We, as humans, rely on energy for every aspect in our life, and this trend increases with the civilization development. Energy can be produced from fossil fuels or from renewable energies. Unfortunately, the energy production is mainly made from fossil fuels, which result in the emission of carbon dioxide (CO2). This CO2 has a greenhouse effect that drastically modifies the Earth's climate, resulting in what nowadays we know as the "global warming" effect. The global warming is in particular causing sea levels to rise, resulting in flooding and erosion of coastal and low-lying areas.

For this purpose, renewable energies have captured much of the attention; they can produce energy from clean (and renewable) sources. In fact, for many years renewable energies were the only source of energy. 1 Nonetheless, it was with the discovery of cheaper energy sources such as fossil fuels or nuclear energy, renewable energies were put aside. CO2 emissions are growing since the industrial revolution, with fossil fuels becoming the 66.8% of the total electricity production in 2017. The main countries contributing to the global warming are China (28% of the total fossil fuel emissions), United States (14%) and India (7%). 2 Now we start to see the consequences and renewable energies are gaining interest again. The main sources of renewable energies are wind, hydroelectric, geothermal, tides and sunlight. While wind and hydroelectric sources are the main renewable technologies generating power, the power generation capacity from solar photovoltaics (solar PV) is expected to increase tremendously by 2040, as in Fig. 1.1. Each renewable technology has its own challenges, however, with the technology development of the last years, they are being improved and, at this moment, they are one of the major solutions to global warming. 2 Photovoltaic technologies emerged as one of the simplest energy solutions to implement in order to reduce the effects of the climate change. Solar panels can transform light into energy and can be installed on buildings. They require less infrastructure than other renewable technologies. This makes them a suitable solution for developing countries, especially in regions like Africa where, for instance, wind or hydroelectric technologies are not feasible.

The actual challenge is to substitute fossil fuels, which were providing around 80% of the total energy demand worldwide in 2013. 2 This translates into around 18 TW continuous power in a year. It may seem a huge amount of energy to cover, however, the Earth receives around 0.8 x 10 5 TW of solar power, which could be a sufficient amount in just one hour to supply the whole population for a year. The European photovoltaic industry association represented well the amount of energy that the Earth receives and the capacity of the actual energy technologies, Fig. 1.2. 3

Figure 1.2. Comparison of the capacity of the actual energy technologies with the amount

of energy that the earth receives from sun. 3 Nonetheless, nowadays only 23.7% of the energy consumption is coming from renewable energies. From the renewable energies, solar PV contributes just with a 2.2% but as previously mentioned, their capacity of energy generation is expected to increase to more than 1 TW by 2023, Fig. 1.3. 4 Solar PV gained attention since Becquerel discovered the photovoltaic effect in 1839, but it was in 1954 when solar cells stood out. 5 The Bell Laboratories provided a silicon-based solar cell for a space program, with a power conversion efficiency (PCE) of 4.5%. 6 The main characteristic of silicon solar cells consist in a p-silicon and an n-silicon region, forming a p-n junction. The success of that technology depends on the development of materials that allow high power at a minimum cost. The price has been decreasing for the last 30 years, what makes it even harder for renewable energies to be competitive. Even if silicon-based solar cells were traditionally the best performing solar cells with the best silicon-based solar cell now performing at 26.7% efficiency, they have several constraints. 7 Silicon is made through a highly demanding energy process and, together with their lifetime and costs from the solar panels, make it a technology far from ideal.

However, the lower cost of multi-crystalline silicon panels, compared to crystalline devices, made them the preferred source of energy even if their efficiency was slightly lower, with their highest PCE at 23.3%. 8 Regarding their lifetime, silicon solar cells last around 30 years, and their payback time is meant to be around 10 years, what means only 20 years of effective production in a solar plant. Indeed, the expected amount of energy required for the entire world for 2050 is of 30 TW, in order to supply a similar amount of energy just with solar plants (10 km 2 for 1 GW), an area as large as India would be needed. 4 Great advances have been achieved in making them cheaper and more efficient.

This allowed the installation of solar panels in conventional houses; however, there is still room for improvement. Pollution could be tremendously reduced if more people had access to solar panels. 4 With the purpose to overcome these problems, second generation solar cells were developed. Here, solar cells were fabricated with glass or stainless steel as a substrate and thin films were deposited on top (~ 1 µm). The thickness reduction compared to silicon solar cells (~ 100 µm) allowed for a reduction in cost and weight. The main thin film materials used in the second generation include copper-indium-gallium-diselenide (CIGS), cadmium telluride (CdTe), amorphous silicon (a-Si) copper-zinc-tin-sulfideselenide (CZTSSe). Thin films such as CdTe has a lower efficiency than silicon solar cells, 22.1%, but their advantage lies in their temperature coefficient. 8 CdTe can stand high temperatures better than silicon; at 65 ºC their efficiency loss is around 8.4 % while the loss for silicon solar cells is of around 18%. 9 CIGS showed high efficiencies in the order of 22-23 %, however, it is necessary to find non-toxic and low cost alternatives for these materials. [10][11][12][13][14] One of the best candidates in second generation solar cells with nontoxic materials is Cu2ZnSnS4 (CZTS). 15 With further efforts to reduce costs and to be competitive against more conventional energy sources, third generation solar cells emerged. This last wave of new technologies gained much attention as their improvements have been done in a short period of time.

Materials used in third generation are not used in a conventional p-n junction system to extract charge carriers anymore, instead they have the ability separate carriers by themselves. They are rather sandwiched between a p and an n material, creating structures such as: p-i-n or n-i-p (they will be further described in the next sections). They can be deposited on flexible surfaces and they are cheaper than the later solar cell generations.

Third generation technologies involve dye-sensitized solar cells (DSSC), quantum dots (QD), organic photovoltaic (OPV) and, lately, perovskite solar cells (PSC). Dyesensitized solar cells were the most popular since the early 90s. Among the typical dyes were ruthenium adsorbed to a mesoporous layer of titanium dioxide (TiO2) or cobalt (III/II) redox complexes. DSSC efficiency is now around 15%, which is a low efficiency when compared to later generations. 16 The introduction of perovskites as a DSSC did not gain much interest but, when Chung et al. used them in solid state (PSC), they increased their efficiency dramatically up to 10.2% and its popularity increased as well. 17 They are now one of the most promising alternatives to silicon solar cells. In this thesis, we focus in perovskite materials, one of the most popular materials not only for third generation PV technologies but also in the energy field.

Perovskite Properties

Perovskites have been found to make highly efficient solar cells. Recent reports have demonstrated that one day they could replace silicon as in their very rapid development, they have reached efficiencies of 25.2%, a level reached by silicon-based solar cells only 20 years ago. 19 Because of their long charge diffusion lengths and low defect densities they have low recombination rates, thus, perovskites enabled a breakthrough in solution processed photovoltaic devices. Besides, perovskite solar cells are manufactured using simple upscalable solution processing methods. As such, this technology presents a really serious alternative to silicon solar cell technology. The prospects seem to be as bright for other optoelectronic applications of the perovskite materials, for example in display technologies. 20,21

Crystal structure

Perovskites typically have an ABX3 crystal structure, Fig. 1.5, composed of five atoms per unit cell, with metal cation B at the center, cation A at the corners, and six anions X shaping the corners of an octahedron and placed at the center of six planes of the unit cell. Perovskite structures can vary depending on the cation in the A site. For instance, CH3NH3PbI3 (MAPbI3) can adopt different structures upon heating, see Table 1. 22,23 MAPbI3 perovskite is orthorhombic below 161 K, tetragonal from 161 to 330 K and cubic above 330 K. 22 In fact, substituting methylammonium (CH3NH3, MA) cation with bigger cations such as formamidinium (CH(NH2)2 + , FA), ethylammonium (CH3CH2NH3 + , EA), butylammonium (CH3(CH2)3NH3 + , BA), guanidinium (CH3(NH2)3 + , GA) or combination of cations, can lead to different structures or even form a layered structure instead of a 3D one. The most commonly used ratio to predict the perovskite formation is the Goldschmidt tolerance factor (t), defined as 24

𝑡 =

𝑟 𝐴 𝑟 𝑥 √2(𝑟 𝐵 +𝑟 𝑥 ) (1.1) Where rA and rB are the ionic radius of the A and B cations respectively and, rx is the ionic radius of the anion. The tolerance factor is used to assess if the A site cation fits in the cavities of the BX3 framework. Perovskites form between 0.8 and 1, while cubic perovskites form within the range 0.9-1.0, distorted perovskites can be found between 0.8-0.9 due to tilting of the BX6 octahedra and lowering the symmetry. If we go out of the range say lower than 0.8, the perovskite may be too distorted, if we go around a tolerance factor of 1.1, the A site cation is too large and thus the perovskite cannot be formed. A second parameter used together with the tolerance factor is the octahedral factor (µ), this factor predicts the stability of the BX6 -4 octhaedral framework, it is given by The octahedral factor values range from 0.44 to 0.90 for a stable structure. 26 It is known that analogues of MAPbI3 made progress in photovoltaic applications and they were introduced in other optoelectronic applications, however, they are moisture sensitive and have low thermal stability. 27 Other A cations showed an improved stability, carrier lifetime and diffusion length in the form of single crystal and solution processed films. In addition, FA + has a larger ion radius than MA + , which leads to a higher tolerance factor that is closer to 1. 28 Cation A can enhance the properties and, a combination of cations could make improvements that are even more significant. Anion X usually is bromine, iodine or chlorine.

µ = 𝑟 𝐵 𝑟 𝑥 (1.2)

Optoelectronic Properties of Perovskites

Organic-inorganic halide perovskites are very interesting for photovoltaic applications due to their optoelectronic properties. They show a very good optical absorption due to their tunable bandgap (see Fig. 1.6a and b), long diffusion lengths, high carrier mobility and high tolerance to defects. [29][30][31][32][33][34][35][36] In particular, MAPbI3 is the most studied perovskite material in solar cells and its properties are believed to be representative for most of the perovskites. MAPbI3 shows a sharp absorption peak, indicating a direct bandgap of 1.55 eV. [29][30][31][32][33][34] Indeed, the MAPbI3 absorption coefficient is so high that the absorption in the visible spectrum is still achievable at thicknesses in the order of 300 nm, which is much thinner than the typical thickness (above 1 µm) required for other solar cell active layers.

The absorption coefficient below the bandgap shows the Urbach tail or Urbach energy. 37 The Urbach energy determines the degree of structural and electronic disorder within the material. 40,41 This effect was detected when measuring the absorption of a defective perovskite crystal, defects induce additional states extended a few meV from the sharp band edges, broadening the Urbach tail. 42 The Urbach value obtained for MAPbI3 is as low as 15 meV, explaining the high Voc in perovskite solar cells. 43 Another property in perovskite solar cells is their high charge mobility, around 25 cm 2 /V•s, related to the ionicity of the metal-halide bond. 44 A high mobility and a low non-radiative recombination leads to carrier diffusion lengths are as large as 100 nm for both electrons and holes in MAPbI3. it increases up to 175 nm in MAPbI3 single crystals due to the reduction of defects. 45 All the above described properties and findings together with a PSC device structure optimization allowed for efficiencies as high as 24.2%. 46 With efficiencies approaching the theoretical maximum efficiency (30.14%), also known as the Shockley-Queisser (S-Q) limit, the question that arises is how to continue reducing the gap. 47 If we take as a reference the reported PSC with 24.2% efficiency, the FF x Voc product is still far from the theoretical value by 25%. 46 This means that research must be directed towards the origin of the FF and Voc lowering effects. So far, the causing effect has been addressed to non-radiative charge recombination and it can be divided in two processes: Shockley-Read-Hall (SRH) recombination and interface recombination. 48 Another typical recombination process is band-to-band recombination but this one is radiative recombination and when there is only radiative recombination, the PSCs efficiencies could theoretically reach the S-Q limit. 47 However, non-radiative recombination processes are still present in perovskites and their origin is mainly due to defects in the crystal structure. Moreover, defects are not only responsible of lower PSC efficiencies but also they are detrimental for their stability. All these detrimental effects induce hysteresis during I-V scans.

Working Mechanism of Perovskite Solar Cells

A perovskite solar cell is, in general terms, composed of the perovskite material sandwiched between two electrodes, one of which is transparent to allow the light to pass through and reach the perovskite layer. This layer is commonly covered by a hole transport layer and an electron transport layer, HTL and ETL respectively. Multiple structures have been developed to improve the PCE, however, they can be classified in two major groups: n-i-p structures, where the perovskite is deposited on top of the ETL as in Fig. 1.7a; and p-i-n structures, where the perovskite is deposited on top of the HTL as in Fig. 1.7b. In n-i-p structures, the most common structure contains a mesoporous TiO2 layer as ETL. Since the contact area with the photoactive layer was increased, it was believed that it helped with the extraction of charges. Later on, other structures without the mesoporous scaffold showed high efficiencies as well, proving that charges can be transported to the respective electrodes by the perovskite itself. 49 Solar cells working principle relies on the absorption of light (photons). The absorption of light produces electrical energy in the form of current and voltage. For that purpose, the photoactive material in the solar cell absorbs the light when the photon energy is higher than the bandgap of the material. Therefore, the bandgap of the active material should be minimized. We expect the efficiency of the solar cell to increase with the light absorption capability, however, when the bandgap energy is much lower than the photon energy, there is an excess of energy which is lost. Bandgap optimization is then a tradeoff between photon harvesting and maximization of the energy of extracted charges.

To evaluate the performance of terrestrial solar cells, the light source irradiated into the solar cell is standardized. The irradiation should have an air mass 1.5 global (AM1.5G), with an intensity of 1000 W/m 2 . 54 The AM1.5G includes direct and diffuse light only.

With a standard procedure and with a solar simulator it is easy to compare the performance of solar cells regardless of the weather.

The light absorption promotes an electron from the ground state to an excited state. The highest energy where the electrons can be found is the valence band (VB) while the lowest energy band where there are not electrons is called the conduction band (CB). The transition of an electron to an excited state in the CB creates an electron (negative charge) and a hole (positive charge) carrier pair, as represented in Fig. 1.8. The electron-hole pair (also called an exciton) attracts each other with a binding energy (Eb) of the order of < 5 meV for perovskites. 55 Since the Eb is lower than the thermal energy at room temperature (around 25 meV), the light absorption will create free charges. 56

𝑛 = ∫ 𝑁(𝐸) 𝐹(𝐸) 𝑑𝐸 𝐸 𝑡𝑜𝑝 𝐸 𝑐 ≅ 𝑁 𝐶 𝑒 (𝜂 𝑒 -𝐸 𝐶 ) 𝑘 𝐵 𝑇 ⁄ (1.3)
where N(E) is the density of states and F(E) is the Fermi-Dirac distribution. 57 Then, following this equation, the electrochemical potential (ηe) of an electron in the CB is

𝜂 𝑒 = 𝐸 𝑐 -𝑘 𝐵 𝑇 𝑙𝑛 𝑁 𝑐 𝑛 (1.4)
where kB is the Boltzmann constant, n is the electron density, Nc is the effective density of states (amount available states close to the band edge) and T the temperature. If we then take into account that the electrochemical potentials are quasi-Fermi energies that may be different between the different layers when an external bias V is applied, the concentration of electrons (n) and holes (p) in the conduction band and valence band respectively is given by

𝑛 = 𝑁 𝑐 exp (- 𝐸 𝑐 -𝐸 𝐹 𝑛 𝑘 𝐵 𝑇 ) (1.5) 𝑝 = 𝑁 𝑣 exp (- 𝐸 𝐹 𝑝 -𝐸 𝑣 𝑘 𝐵 𝑇 ) (1.6)
where EF n is the Fermi level of n-type semiconductors and EF p is the Fermi level of ptype semiconductors. From Eq. 1.5 and Eq. 1.6 and taking into account that the charge of an electron times the photovoltage follows the relationship: qV = EF p -EF n , we obtain

𝑛𝑝 = 𝑁 𝑐 𝑁 𝑣 exp (- 𝐸 𝑔 -𝑞𝑉 𝑘 𝐵 𝑇 ) (1.7)
where Eg is the bandgap energy of the perovskite. The probability of an electron to find a hole and recombine is higher in open-circuit condition, as previously discussed, then

𝑞𝑉 𝑜𝑐 = 𝐸 𝑔 -𝑘 𝐵 T ln 𝑁 𝐶 𝑁 𝑉 𝑛𝑝 (1.8)
where there is a relationship between Voc, np product, temperature and bandgap.

The short-circuit current density (Jsc) is the maximum photo-generated current that can be achieved. It is the point where no voltage is applied between the anode and the cathode (see Fig. 1.9). In fact, for an ideal diode, the photocurrent is

𝐽 𝑖𝑑𝑒𝑎𝑙 = 𝐽 𝑠𝑐 -𝐽 0 [exp ( 𝑞𝑉 𝑘 𝐵 𝑇 ) -1] (1.9)
where J0 is the saturation current. Now, the ideal open-circuit voltage when only radiative recombination is present, can be defined as 58

𝑉 𝑜𝑐 = 𝑘 𝐵 𝑇 𝑞 𝑙𝑛 ( 𝐽 𝑠𝑐 𝐽 0 + 1) (1.10)
The maximum electrical power (Pmax) point is the position on the I-V curve where the maximum product between current density (Jmpp) and voltage (Vmpp) is reached. The fill factor (FF) describes the ideality of a solar cell, that is, how close is the J-V curve to form a rectangle like in Figure 1.9 and can be expressed as

𝐹𝐹 = 𝐽 𝑚𝑝𝑝 𝑉 𝑚𝑝𝑝 𝐽 𝑠𝑐 𝑉 𝑜𝑐 (1.11)
The power conversion efficiency (PCE) determines the performance of a solar cell. It is the ratio between the output and input power, Pout and Pin respectively.

𝑃𝐶𝐸 = 𝑃 𝑜𝑢𝑡 𝑃 𝑖𝑛 (1.10)
Still, the PCE can be hindered by multiple reasons. For perovskite solar cells two of the main issues are the hysteresis in the I-V measurements and the presence of defects within the material.

Hysteresis

Hysteresis in perovskite solar cells refers to the divergence between the I-V curves measured in forward (from short circuit to open circuit) and reverse (from open circuit to short circuit) scan directions and the time scale for this effect is ~ 100 sec. 59,60 Hysteresis has a temporal-dependence behaviour, what means that can be modified with parameters such as scan speed, scan direction or scan range. 61 Altogether, it is an evidence of poor stability. 62,63 A typical hysteretic behaviour was previously shown in Fig. 1.10 and the general trend is to measure a higher current in reverse scan direction than in forward scan direction. The inset table gives the determined performance parameters. 62 This means that one or more processes of the solar cell working mechanism (before described) is being disturbed. Considering that carrier generation and carrier transport are rather fast processes, the lower carrier collection seen in forward scan must be attributed to the transport at the interfaces. Although the origin of the hysteresis effect is still a subject of debate, three main approaches have been proposed to be causing it: ferroelectricity, charge trapping/detrapping and ion migration.

Ferroelectricity

Ferroelectricity appears when spontaneous polarization (P), generated from noncentrosymmetric structures, can be switched by an external electric field. Several groups claimed that MAPbI3 exhibits a ferroelectric behaviour despite being a centrosymmetric tetragonal space group of I4/mcm. [START_REF] Stoumpos | Semiconducting Tin and Lead Iodide Perovskites with Organic Cations: Phase Transitions, High Mobilities, and Near-Infrared Photoluminescent Properties[END_REF][START_REF] Quarti | Interplay of Orientational Order and Electronic Structure in Methylammonium Lead Iodide: Implications for Solar Cell Operation[END_REF] Nonetheless, the central MA cation is mobile and can change its orientation, thus decreasing the lattice symmetry. [START_REF] Bakulin | Real-Time Observation of Organic Cation Reorientation in Methylammonium Lead Iodide Perovskites[END_REF] Several ferroelectric domains were detected using Piezoresponse force microscopy (PFM), a tool used for probing local ferroelectric response. [START_REF] Stoumpos | Semiconducting Tin and Lead Iodide Perovskites with Organic Cations: Phase Transitions, High Mobilities, and Near-Infrared Photoluminescent Properties[END_REF][START_REF] Kutes | Direct Observation of Ferroelectric Domains in Solution-Processed CH3NH3PbI3 Perovskite Thin Films[END_REF][START_REF] Kim | Ferroelectric Polarization in CH3NH3PbI3 Perovskite[END_REF][START_REF] Chen | Interface band structure engineering by ferroelectric polarization in perovskite solar cells[END_REF][START_REF] Coll | Polarization Switching and Light-Enhanced Piezoelectricity in Lead Halide Perovskites[END_REF] The polarization of these domains alters the band structure at the interface, resulting in a modified I-V performance between forward and reverse scans. Ferroelectric polarization was demonstrated by Wei et al., who proposed a model for this behaviour. [START_REF] Wei | Hysteresis Analysis Based on the Ferroelectric Effect in Hybrid Perovskite Solar Cells[END_REF] With an external field, the MA cation can align to follow the direction of the applied field, resulting in a depolarization of the field.

However, P cannot follow the speed of the voltage scan, so the excess polarization generates a transient current state. This excess polarization on forward scan direction is parallel to the internal field, what makes the transient current to be lower than in reverse scan direction. Nonetheless, models such as the one described for the perovskite ferroelectricity are questioned by number of other groups. [START_REF] Kutes | Direct Observation of Ferroelectric Domains in Solution-Processed CH3NH3PbI3 Perovskite Thin Films[END_REF][START_REF] Leguy | The dynamics of methylammonium ions in hybrid organic-inorganic perovskite solar cells[END_REF][START_REF] Fan | Ferroelectricity of CH3NH3PbI3 Perovskite[END_REF] Some claim that there is no ferroelectric behaviour in perovskites and some other pointed that even if Wei et al.

found ferroelectric properties, they are too small to cause the hysteretic behaviour seen by other groups. [START_REF] Wei | Hysteresis Analysis Based on the Ferroelectric Effect in Hybrid Perovskite Solar Cells[END_REF]

Ion Migration

Ion migration being the cause of J-V hysteresis gained more attention than ferroelectricity. [START_REF] Richardson | Can slow-moving ions explain hysteresis in the current-voltage curves of perovskite solar cells?[END_REF] Here, under an external electrical field, positive and negative ions are driven towards the respective electrodes. Ions accumulate at the interface between the transport layers and the perovskite, resulting in a change of the polarity and the modulation of the interfacial barrier energies, thus producing hysteresis. [START_REF] Zhang | Charge selective contacts, mobile ions and anomalous hysteresis in organicinorganic perovskite solar cells[END_REF][START_REF] Zhao | Anomalously large interface charge in polarity-switchable photovoltaic devices: an indication of mobile ions in organic-inorganic halide perovskites[END_REF] In a perovskite system, the elements that can diffuse are the organic cations, halide anions and metallic ions. However, heavy metal cations like Pb +2 have low mobility and are not considered as a diffusive part. Smaller organic cations will diffuse faster than larger ones and, halides are considered to have the highest mobility compared to its counterparts. [START_REF] Haruyama | First-Principles Study of Ion Diffusion in Perovskite Solar Cell Sensitizers[END_REF][START_REF] Eames | Ionic transport in hybrid lead iodide perovskite solar cells[END_REF] Regarding the diffusion of organic cations like MA + , it was observed an increase in concentration when poling the sample. [START_REF] Yuan | Photovoltaic Switching Mechanism in Lateral Structure Hybrid Perovskite Solar Cells[END_REF] They accumulate at the interface and align with the field, hindering charge extraction. [START_REF] Azpiroz | Defect migration in methylammonium lead iodide and its role in perovskite solar cell operation[END_REF] However, the estimated diffusion coefficient of iodide ions in MAPbI3 ranges from 10 -11 to 10 -12 cm 2 /s, far from the estimated value for MA + cations, 10 -16 cm 2 /s. [START_REF] Eames | Ionic transport in hybrid lead iodide perovskite solar cells[END_REF] Meanwhile, Li et al. detected the migration of iodide anions to the positive electrode, which have an estimated coefficient of diffusion within the range before mentioned. [START_REF] Li | Iodine Migration and its Effect on Hysteresis in Perovskite Solar Cells[END_REF] The migration of iodide ions can generate hydrogen ions, impacting the mobile defects. [START_REF] Egger | Theory of Hydrogen Migration in Organic-Inorganic Halide Perovskites[END_REF] However, these hydrogen ions are not considered to be responsible of the observed hysteresis since their concentration remains low due to the weak acidity of methylammonium. [START_REF] Frost | What Is Moving in Hybrid Halide Perovskite Solar Cells?[END_REF] When iodide ions migrate, they leave vacancies, which are driven towards the negative electrode and, Xiao et al. suggested they are responsible for n-doping under an external electric field. [START_REF] Xiao | Giant switchable photovoltaic effect in organometal trihalide perovskite devices[END_REF] The ion migration was proposed to take place by a hopping mechanism. Ions will migrate through vacant sites: MA + goes into a vacant cage, and iodide ions move along the Pb-I octahedra. [START_REF] Eames | Ionic transport in hybrid lead iodide perovskite solar cells[END_REF] Migration can occur when the PbX3 framework opens, and the energy required is called the activation energy (Ea). Therefore, iodide ions can migrate through iodide vacancies (VI) or iodide interstitials (Ii), with a lower activation energy required for migration through vacancies. [START_REF] Yu | Native Defect-Induced Hysteresis Behavior in Organolead Iodide Perovskite Solar Cells[END_REF] Therefore, it seems reasonable that hysteresis can be originated from ion migration. However, other groups investigated hysteresis effects with the same MAPbI3 crystal but with other ETLs and HTLs, establishing lower hysteresis features. [START_REF] Heo | Hysteresis-less inverted CH3NH3PbI3 planar perovskite hybrid solar cells with 18.1% power conversion efficiency[END_REF] This behaviour means that ion migration cannot be responsible for the hysteresis effect alone. Instead, a combination of both, ion migration and trap states might be a better explanation. [START_REF] Richardson | Can slow-moving ions explain hysteresis in the current-voltage curves of perovskite solar cells?[END_REF][START_REF] Yu | Native Defect-Induced Hysteresis Behavior in Organolead Iodide Perovskite Solar Cells[END_REF][START_REF] Van Reenen | Modeling Anomalous Hysteresis in Perovskite Solar Cells[END_REF]

Charge trapping/detrapping

Another factor that should be taken into account is the relationship of trap states with hysteresis effects. Trap states have also been proposed as the origin of hysteresis since the existence of traps in perovskites has been widely discussed. [START_REF] Van Reenen | Modeling Anomalous Hysteresis in Perovskite Solar Cells[END_REF][START_REF] Yuan | Perovskite Solar Cells: Hot-Electron Injection in a Sandwiched TiOx-Au-TiOx Structure for High-Performance Planar Perovskite Solar Cells[END_REF] Trap states may appear with low temperature fabrication processes (solution process) and, ultimately, they affect charge transport, separation and recombination. [START_REF] Yuan | Perovskite Solar Cells: Hot-Electron Injection in a Sandwiched TiOx-Au-TiOx Structure for High-Performance Planar Perovskite Solar Cells[END_REF] Among the possible traps, shallow and deep, it is true that perovskites show low Urbach energies, in the order of 20-40 meV, which implies a high structural quality and electronic properties together with high mobilities. [START_REF] Singh | Effect of Thermal and Structural Disorder on the Electronic Structure of Hybrid Perovskite Semiconductor CH3NH3PbI3[END_REF] On the contrary, deep traps are associated to Shockley-Read-Hall nonradiative recombination and lead to a performance losses. [START_REF] Miller | Intense Internal and External Fluorescence as Solar Cells Approach the Shockley-Queisser Efficiency Limit[END_REF][START_REF] Wetzelaer | Trap-Assisted Non-Radiative Recombination in Organic-Inorganic Perovskite Solar Cells[END_REF] A mechanism of trappingdetrapping was described to produce hysteresis. [START_REF] Li | Hysteresis mechanism in perovskite photovoltaic devices and its potential application for multi-bit memory devices[END_REF] Shallow traps can be filled by charges and then released with changing bias while deep trap states can produce a change in the electric field, resulting in a hysteresis in J-V performance. [START_REF] Li | Hysteresis mechanism in perovskite photovoltaic devices and its potential application for multi-bit memory devices[END_REF] Even if perovskites are found to have low trap densities, they are dominantly accumulated at the interface or at the surface where the structure is not as crystalline as the bulk. 56,[START_REF] Yamada | Photocarrier Recombination Dynamics in Perovskite CH3NH3PbI3 for Solar Cell Applications[END_REF] However, these traps can also be driven to the surface due to strong electron-phonon coupling. [START_REF] Wu | Trap States in Lead Iodide Perovskites[END_REF] The accumulation phenomena has been observed by Kelvin Probe Force Microscopy (KPFM) at the interface between the perovskite and the HTL after turning on the light. [START_REF] Bergmann | Local Time-Dependent Charging in a Perovskite Solar Cell[END_REF][START_REF] Bergmann | Real-space observation of unbalanced charge distribution inside a perovskitesensitized solar cell[END_REF] These traps can be passivated by adding additional thin films of fullerene, which was observed to reduce hysteresis. [START_REF] Xing | Interfacial Electron Transfer Barrier at Compact TiO2/CH3NH3PbI3 Heterojunction[END_REF][START_REF] Shao | Origin and elimination of photocurrent hysteresis by fullerene passivation in CH 3 NH 3 PbI 3 planar heterojunction solar cells[END_REF] Even if these findings suggest that traps are behind the hysteresis effect, the trapping and detrapping effect takes place in the order of ms or ns, which is far from the hysteresis times. The lower trap density in perovskites makes it also hard to believe that hysteresis is only related to the traps. [START_REF] Oga | Improved Understanding of the Electronic and Energetic Landscapes of Perovskite Solar Cells: High Local Charge Carrier Mobility, Reduced Recombination, and Extremely Shallow Traps[END_REF]100 Finally, traps should be filled with increasing illumination intensity, thus decreasing the hysteresis, what happens to be the contrary to what it is observed. As a matter of fact, the actual debate is not whether hysteresis arises from trap states or ion migration. It is about how much the hysteresis is affected from ion migration and how much from trap states.

Structural Defects in Perovskites

As previously discussed, defects can create traps and are detrimental for PSC performance, however, there is a big discussion about the role they play in the PSC stability. To understand how defects affect the electronic structure of perovskite materials, the origin of the electronic levels must be clarified.

Energy band structure

In MAPbI3, the electronic levels of the MA cation lie deep within the valence and conduction bands (VB and CB respectively), which means that the band edges are determined by the BX6 octahedra. Due to the ionicity between the Pb and I atoms in MAPbI3, the CB is determined from the Pb p orbitals. The covalent antibonding coupling between Pb p and I p states is not strong, therefore, the CB energy is just slightly higher than the atomic Pb p states, as in Traps will be filled, thus inducing higher Voc. 102

Structural defects

Defects are normally addressed in the bulk, however defects present at the surface cause electronic heterogeneities and condition the electronic properties of the perovskite, which are directly involved in the charge extraction. 104 Traps can be classified as shallow or deep, which means they are located either close to the band-edges or close to the midgap, respectively. Specifically, the possible point defects in MAPbI3 can be induced by Pb or MA vacancies (VPb and VMA, respectively), substitution of I atoms at Pb or at MA sites (IPb and IMA, respectively), substitution of MA ions at Pb sites (MAPb) and I interstitials (Ii). Calculations on the formation energies of all the possible point defects can be seen in Fig. 1.12 and revealed that shallow trap states are much more likely than deep traps, and they can be an explanation for the long electron-hole diffusion lengths observed in these materials. All these studies are also pointing out how the type and density of traps at the surface of the perovskite are depending on the growth conditions of the materials along with the anisotropy of the surface properties. In spite of the type of defects present at the surface, their intrinsic role against stability is still under debate. To enhance perovskite stability, it is crucial to understand how these defects affect the bulk properties and how they interact with water molecules, especially at the early stage of degradation. Together with defects, other factors are also causing some concerns in perovskite solar cells. The toxicity of lead, their I-V hysteresis losses and their environmental stability makes them a potential technology to be further improved.

Stability of Perovskite Solar Cells

At atmospheric conditions, perovskite materials suffer degradation from a number of environmental factors. They can be classified in three groups: degradation induced by the illumination, the temperature and/or the moisture. A wide range of stability studies have been performed for each degradation process, however, there is mainly one thing they have in common: the majority of them were carried out using MAPbI3. It is the most studied perovskite so far and the perovskite studied in this work. Therefore, here there are some of the most remarkable findings concerning stability of PSC.

Prolonged radiation exposure of PSCs, specially UV radiation, is one of the characteristic degradation processes seen in perovskites. 108 Radiation instabilities were first seen at the interface between the mesoporous titanium oxide (TiO2) and the perovskite. Oxygen vacancies or Ti 3+ exist in the deep level electron sites of TiO2, which facilitate the interaction with the oxygen creating a charge transfer complex. 109 After light absorption, an electron-hole pair is created and separated to the CB and VB respectively. The hole at the VB recombines with the electron from the oxygen at the adsorption site, creating an excess of electrons and, at the same time, creating oxygen vacancies in the TiO2 surface. [111][112][113] To overcome this problem, one of the most popular solution was to introduce mesoporous Al2O3 to replace the mesoporous layer of TiO2 as in Fig. 1.13. 114-116 Fig. 1. 13 shows the work where the Al2O3 mesoporous layer was compared to the TiO2 mesoporous layer, showing a better electron blocking. 110 The perovskite suffers also photo-degradation by itself, and even more when it is exposed to oxygen at the same time. 117 Regarding the effects of UV-light illumination in the perovskite layer, the inorganic PbI 2+ can capture two electrons and result in Pb and I decomposition. [118][119][120] As a solution it was suggested to add an excess of PbI 2 . 121 Altogether, illumination in PSCs may induce a bandgap change and an interfacial geometry change. 122 Temperature conditions are also an important factor for the PSC performance. When combined with light radiation, the MAPbI3 perovskite starts the degradation process at 75 ºC due to the weak Pb-I bond, which can be broken by incident photons. However, without light influence, perovskite decomposition starts above 85 ºC, specially suffering high decomposition at 95 ºC, following the Reaction 1.16. 120,123 The exact degradation mechanism is not well understood, however, the consensus is that perovskites suffer tremendous degradation above 100 ºC. 124 Thermal degradation, contrary to light degradation, starts from the perovskite surface and it continues progressively in the bulk, see Fig. 1.14. 124 The degradation mechanism is still under debate, with a general consensus claiming that high temperature is enough to degrade the perovskite layer.

𝐶𝐻 3 𝑁𝐻 3 𝑃𝑏𝐼 3 (𝑠) 𝛥 → 𝑁𝐻 3 (𝑔) + 𝐶𝐻 3 𝐼 (𝑔) + 𝑃𝑏𝐼 2 (𝑠) (1.16)
Moisture degradation is one of the major concerns in the perovskite community too. 125- 130 With an organic cation (like MA + or FA + ) in the center of the perovskite structure, it makes the perovskite extremely vulnerable to water interactions. That is why, the PSCs are fabricated in inert atmospheres such as a nitrogen-filled glovebox. This type of degradation is not clearer than the ones described before, that is why many groups are trying to solve this critical issue. 33,127,[131][132][133][134][135][136] One of the first degradation mechanisms proposed was the Grotthuss mechanism, here the proton in the MA + cation is transferred to the water molecules. Water molecules transfer the proton to the halide compound, resulting in HI and PbI2 formation. 132 Then, some groups proposed that with the sole presence of water, the perovskite structure was directly degraded into PbI2. 33,137 Most recently, several groups claimed that intermediate hydrated-phases such as CH3NH3PbI2•H2O and/or (CH3NH3)4PbI6•H2O are formed in the first moments of moisture degradation to finally decompose into PbI2, Fig. 1.15. 127,133,134 Interestingly, they were detected by X-ray diffraction studies. 127,133 These studies were more keen to a degradation mechanism starting with perovskite hydration rather than an acidic-base reaction proposed in the beginning (Grotthuss mechanism). More detailed studies by Leguy et al. showed that the hydrated intermediates are reversible upon drying as in the following reaction: 135 bond. 121,125,138,139 As one can imagine, even if the mechanism behind the moistureinduced degradation is still under debate, there is a general consensus about the presence of intermediate phases appearing before irreversible degradation. Now, the possible degradation mechanisms can be minimized through the combination of different transport layers, which not only transport charges but also can protect the perovskite from degradation.

4𝐶𝐻 3 𝑁𝐻 3 𝑃𝑏𝐼 3 + 4𝐻 2 𝑂 ⇆ 4[𝐶𝐻 3 𝑁𝐻 3 𝑃𝑏𝐼 3 • 𝐻 2 𝑂] (1.17) 4[𝐶𝐻 3 𝑁𝐻 3 𝑁𝐻 2 • 𝐻 2 𝑂] ⇆ (𝐶𝐻 3 𝑁𝐻 3 ) 4 + 𝑃𝑏𝐼 6 • 𝐻 2 𝑂 + 3𝑃𝑏𝐼 2 + 2𝐻 2 (1.18)

Charge Transport Layers in Perovskite Solar Cells

Stability issues can get worse or can be reduced depending on the transport layers in which the perovskite is in contact. There are several options for both, the electron transport layer and the hole transport layer. The most evident important characteristic to pay attention when choosing the transport materials is the energy band alignment. 140 However, other aspects such as stability of the material or their degree of solubility will directly affect the solar cell performance. A degraded transport layer will perturb the carrier collection efficiency and, subsequently, the PCE. Transport materials can be either organic or inorganic ones. For the ETL, suitable compounds should be related to negligible hysteresis, flexibility and easy fabrication at low temperatures. [START_REF] Heo | Hysteresis-less inverted CH3NH3PbI3 planar perovskite hybrid solar cells with 18.1% power conversion efficiency[END_REF][START_REF] Shao | Origin and elimination of photocurrent hysteresis by fullerene passivation in CH 3 NH 3 PbI 3 planar heterojunction solar cells[END_REF]141,142 The typical ETLs used in n-i-p structure are TiO2, SnO2 and ZnO. It can be used as a compact layer or as a mesoporous layer but it has been in the mesoporous configuration where they showed higher carrier extraction due to its large area of contact with the perovskite. 143 TiO2 can be deposited in multiple ways. The preferred deposition techniques are spray-pyrolysis, sol-gel, atomic layer deposition and electrodeposition because of the advantage of low temperature processing over other techniques such as radiofrequency sputtering, electron-beam evaporation or thermal oxidation. [144][145][146] However, from section 1.3.4 we know the mesoporous TiO2 layer has a poor stability. In fact the TiO2 layer can lower the PCE by dropping 10-30% of its initial value after several hours. 147 As a result, alternatives such as ZnO emerged, showing higher electron mobility and the ability to be flexible. 148 Moreover, doping the ZnO with nitrogen or aluminium increased its stability and, at the same time, kept a high PCE. 149 SnO2 was another alternative to TiO2 due to its deeper conduction band. As counterparts, its bandgap is wide, which means less UV light absorption.

On the other hand, in p-i-n structure the most common ETL are fullerene derivatives such as [6,6]-phenyl-C61-butyric acid methyl ester (PCBM) or C60, and f respectively. However, it is worth to mention other materials used as HTL in PSC such as CuSCN or NiOx, with high stabilities and performances, however, they are not used as much as PEDOT:PSS or Spiro-OMeTAD. PEDOT:PSS is a mixture of the PEDOT polymer with PSS, which makes the PEDOT soluble. PEDOT:PSS is the most used material as HTL in p-i-n structures due to its ease of fabrication by solution process, transparency and appropriate valence band energy with respect to those of perovskites.

However, since PEDOT:PSS can accelerate the perovskite degradation, MoO3 was introduced to improve its stability. 151,152 Spiro-OMeTAD is the most used material as HTL in n-i-p structures because of energy alignment and amorphous character. 153 However, it is doped with Bis(trifluoromethane)sul-fonimide lithium salt (LiTFSI) , Fig. Regardless of the intrinsic properties of the charge transport layers above mentioned, the growth conditions of the perovskite should be taken into account. Even if the growth process for perovskites is not well understood, it is known that depending on the material beneath the perovskite, the growth conditions will differ, thus, its electronic properties.

Some studies showed a preferential order when the perovskite is deposited on top of TiO2 due to local interactions. 157 There are also studies that compare the energy band positions of the perovskite when deposited on different substrates (for instance TiO2, ZnO and PEDOT:PSS). 158 Therefore, even if the focus of this thesis is not related to PSC performance directly, the choice of a transport layer is also important. Since the tests that were going to be performed to characterize the perovskite layer are slow by themselves, a fast deposition process was preferred. In this aspect PEDOT:PSS was chosen, not only it allows for a relatively fast deposition process but it is also highly reproducible, cheap and very common material used as HTL in perovskite solar cells. Although it was described earlier that it can be unstable, it is sandwiched between the back electrode and the perovskite, limiting its contact with atmospheric conditions, moreover, it is used in controlled environments and under an inert N2 atmosphere.

Objectives and Outline of the Thesis

From the above considerations, it is clear that there are many detrimental effects in perovskite materials causing them to outperform and, lately, to be chemically degraded.

The understanding of the processes taking place in the degradation process is of high importance for their commercialization as solar cell devices. In this regard, the ultimate goal of this work is to contribute to the development of perovskite-based solar cells by providing experimental insight about the perovskite reactivity upon water exposure at the nano-level.

To do so, scanning probe microscopy techniques have been used through all the thesis to understand the electronic behaviour of the perovskite surface which, ultimately, could help to understand the behaviour at the interface between the perovskite and the eventual transport layer in a solar cell stack. After the introduction of this chapter, Chapter 2 provides a description of the sample preparation conditions and the used techniques through this thesis. Chapter 3 compares the local electronic behaviour of two perovskite synthesis routes both leading to MAPbI3. The discussion lies around the presence of two dominating domains at the surface of one of the synthesis routes (MAPI-PbAc2), with a difference in work function between these domains. The main experimental measurements were performed by KPFM and C-AFM.

The optimization of the measuring parameters and conditions has taken quite some work in Chapter 4. This includes not only testing both perovskite synthesis routes by Scanning Microwave Microscopy (SMM) but also finding the optimal procedure to extract some meaningful results. The chapter discusses about the suitability of the mentioned technique to characterize a semiconductor such as perovskite. Continuing with the characterization of the two different perovskite synthesis routes, we use the capabilities of the SMM to track any hysteretic effect.

Chapter 5 takes a closer on the surface stability from both MAPbI3 perovskites upon water exposure. Based on these results, we perform an evaluation of which perturbing phenomena is making the stability difference. In the last part of the thesis, in Chapter 6, final conclusions and outlook of the thesis are presented. Various studies are proposed in view of further complementary analysis of the perovskite surface.

Chapter 2

Materials and Methods

In this thesis, to understand the phenomenon occurring in perovskite solar cells at the nanoscale, a series of preparation methods and experimental techniques have been used.

In order to avoid repetition in the next chapters, the operational principle of the techniques used for characterization will be included in this chapter. Also, the material synthesis and device preparation is described.

Fabrication Process

The device structure prepared in this work is based on the typical materials used for p-in perovskite solar cells, as discussed in section 1. To describe how the perovskite films were made, it should be noted that hybrid perovskites can be fabricated and deposited in multiple ways. However, due to the low cost and ease of use of solution processing techniques, they are the preferred deposition In one-step deposition, a single solution is used so the solvent(s) should dissolve all the perovskite precursors, organic and inorganic, Fig. 2.2a. 1 One-step can become difficult to make for longer organic chains in the perovskite since it would be more difficult to find a good solvent. Same goes for the inorganic part, they are soluble (or stable) with orthogonal solvents. Nonetheless, an optimization of the precursor ratio and the usage of additives makes one-step deposition still a very valid deposition technique. Two-step deposition involves dissolving each precursor individually and spin-coating them sequentially without stopping the spinning, and dropping the second solution at some point of the spin-coating process, Fig. 2.2b. It allows for a better control of the final thickness, morphology and it shows high uniformty. 2 After the deposition, the films are then annealed. Other techniques used in literature but in a lesser extent are vapour process and hybrid processes. In vapour process deposition, perovskites are deposited via evaporation and annealing (as in Fig. 2.2c or 2.2d). 9 This technique is usually done in vacuum and it is especially useful for growing oriented thin films. However, it is not as cheap as solution process and it is limited when depositing different perovskites. The socalled "hybrid" processes are a combination of both, solution and vapour processes. 8 CH3NH3PbI3 perovskite can be formed by different chemical reactions depending on the precursor materials used. The most popular precursors are methylammonium iodide (MAI) and a lead source, typically lead iodide (PbI2), Eq. 2.1. Nonetheless, the perovskite formation can be achieved with other anions as a lead source, or even with a mixture of anions. 10 Zhang et al. studied the role of different anions in the lead precursor and found that they affect the crystal growth kinetics, film morphology and device performance. 11 The typical anion for the PbX2 precursor is iodide. However, lead acetate (PbAc2) was found to accelerate the crystal growth, Eq. 2.1. In addition, the perovskite films can be either p-type or n-type depending on the ratio between the precursors, as in Fig. 2.3. With a PbI2-rich (PbI2-deficient) film, the perovskite will be n-doped (p-doped), respectively. 12 in grey colour are less-likely to form in the films. 12 In this work, two routes were carried out to synthesize MAPbI3. The first one (MAPI-PbAc2 route, described below) was developed by Zhang et al. to increase the grain size and to grow the perovskite with less defects. 13 The second one (MAPI-PbI2 route, described below) is the most used synthesis route for MAPbI3 perovskite and served as a reference route. 14 The purpose is to analyse the electronic differences in the MAPbI3 behaviour (especially at the surface) when using different precursor materials and/or process techniques.

𝑃𝑏𝐴𝑐 2 + 3 𝐶𝐻 3 𝑁𝐻 3 𝐼 ⟶ 𝐶𝐻 3 𝑁𝐻 3 𝑃𝑏𝐼 3 + 2 𝐶𝐻 3 𝑁𝐻 3 𝐴𝑐 (2.1) 𝑃𝑏𝐼 2 + 𝐶𝐻 3 𝑁𝐻 3 𝐼 ⟶ 𝐶𝐻 3 𝑁𝐻 3 𝑃𝑏𝐼 3
Solution process was performed in one-step for the two CH3NH3PbI3 synthesis routes, the materials required for their synthesis are summarized in Table 1. While one route involves the usage of an additive in the precursor solution, the other route is done through an anti-solvent method to improve the morphology. The first synthesis route was made using MAI and PbAc2 as precursor materials and were mixed prior to depositing them in one-step, this perovskite will be referred to as MAPbI3-PbAc2.

MAI (Greatcell Solar) and PbAc2•3H2O (99.999% Merck) were used in a 3:1 molar ratio dissolved with anhydrous N,N-Dimethylformamide (DMF, 99.9% Merck), with a final concentration of 0.5 M. Hydrophosphorous acid (HPA, 50% Merck) was added with a HPA/PbAc2•3H2O molar ratio of 11%. HPA is used to enlarge the grain size and to reduce the non-radiative recombination centres, being 11% the optimal molar ratio. 13 The solution was then spin-coated at 2000 rpm for 60 s, left at room temperature for 10 min.

and heated at 100 ºC for 5 min.

In the second route, MAI and PbI2 were used as precursors and were deposited in onestep using an anti-solvent method, this perovskite will be referred to as MAPbI3-PbI2. 14 The anti-solvent consists of a solvent that speeds up the nucleation via the creation of an instantaneous local supersaturation on the spinning substrate. This solvent is usually dripped after some time while the substrate is spinning, resulting in the formation of an intermediate phase like MAI-PbI2-DMSO. 15 The intermediate phase prevents the fast reaction between MAI and PbI2 precipitation and allows for a uniform morphology. In this thesis, the perovskite was made by mixing lead iodide PbI2 (99% Merck) and MAI in DMF (and 1.5% of the total volume with DMSO as additive) with a 1:1 molar ratio and a final concentration of 1.5 M. The solution was spin coated at 5000 rpm for 30 s while 200 µL of ethyl acetate is dropped into the sample after 10 s to control the end point of the growth process. 14,16 The final thickness for both perovskite routes is of around ~ 300 nm.

Characterization Techniques

X-Ray Diffraction

X-ray diffraction (XRD) is a method used to determine the structural properties of the crystals based on the scattering of light. Since amorphous materials are not organized in a periodic array, they do not scatter light coherently. In a crystal structure each atom is arranged in a periodic array and, when an X-ray interacts with it, a certain amount of energy is absorbed. Electrons occupy certain energy states around an atom, but the absorbed energy is not enough to release the electron, the energy must then be re-emitted in a form of a new x-ray with the same energy as the original, this process is known as elastic scattering. At specific x-ray incident angles, the scattered x-rays can constructively interfere; this is called diffraction and results in a peak in intensity seen in the XRD

𝑛 𝜆 = 2 𝑑 sin 𝜃 (2.2)
where n is an integer known as the order or reflection. The crystal geometry can be determined from the crystal spacing measured by XRD. XRD has been used in this work to determine the crystalline structure of the MAPbI3 perovskite (in Chapter 3) and to determine if any change occur in the structure when exposing the MAPbI3 perovskite to water vapour (in Chapter 5). From the crystal spacing, the lattice distances can be also determined following the expression for a given geometry. Perovskites use to be in a tetragonal configuration as it is the case for the perovskite used for this work:

1 𝑑 ℎ𝑘𝑙 2 = ℎ 2 +𝑘 2 𝑎 2 + 𝑙 2 𝑐 2 (2.3)
X-ray diffractograms were obtained using a Panalytical Empyrean XRD diffractometer operating with the Cu Kα1 (α = 0.1542 nm) source.

X-Ray Photoelectron Spectroscopy

In X-ray photoelectron spectroscopy (XPS), an x-ray is sent to the sample surface with enough energy to eject an electron (photoelectron) as in Fig. 2.5. The threshold energy ν0 at which the electron can be emitted is determined by its frequency. At higher frequencies ν, the emitted electrons will possess higher kinetic energies KEelectron,

𝐾𝐸 𝑒𝑙𝑒𝑐𝑡𝑟𝑜𝑛 = ℎ𝜈 -ℎ𝜐 0 (2.4)
Knowing the KEelectron and the energy of the incident photon, the binding energy (BE) can be calculated as

𝐵𝐸 = ℎ𝜈 -𝐾𝐸 𝑒𝑙𝑒𝑐𝑡𝑟𝑜𝑛 -𝜙 𝑠𝑝𝑒𝑐 (2.5)
where ϕspec is the spectrometer work function. The system is operated under high vacuum (< 10 -9 torr) to avoid surface reactions or contamination. All orbital levels except the s levels are doublets with the two possible states having different binding energies.

UV Visible Spectroscopy

In UV-Visible spectroscopy, a light beam is directed to an absorbing material and then the amount of absorbed (or transmitted) light is determined. The absorbed light depends on the thickness of the material and the concentration following the expression:

𝐴 = 𝛼 𝑐 𝑑 (2.6)
where A is the absorbance, α is the absorption coefficient, c is the concentration of the material and d is the length of the material. The spectrometer shoots light in a range of wavelengths of light, obtaining an absorption spectrum. From the spectrum, the absorption peak at which the material absorbs light can be determined. When the light is absorbed, the electron occupying the valence band is promoted to the conduction band, therefore, a certain amount of energy is absorbed:

𝐸 = ℎ𝜈 = ℎ 𝑐 𝜆 (2.7)
where h is the Planck constant, ν is the frequency of light, c is the speed of light and λ is the wavelength. From this relationship, the optical bandgap of the material can be calculated. However, for materials with a direct allowed transition in the bandgap, the optical bandgap EG could be determined with the Tauc plot (see Fig.

Atomic Force Microscopy -Principle

Atomic Force Microscopy (Multimode AFM from Bruker) is used to provide a topographical image at the nanoscale of the surface under test. AFM probes the forces between the probe and the surface at short distances (between 0 and 10 nm). The scheme of an AFM can be seen in In contact mode (Fig. 2.8a), the probe touches the sample surface with a given force and the probe-surface repulsive forces deflect the cantilever. The force applied to the probe and its deflection are proportional according to Hooke's law (F = -k • δz, where F is the applied force, δz is the deflection distance and k is the spring constant of the cantilever).

Through a feedback loop, the applied force (so the deflection) is maintained. This may be a destructive technique for soft surfaces. In contrast, in non-contact mode (Fig. 2.8b), an ac signal is sent to the piezoelectric material at, or close to, the resonance frequency of the cantilever. The probe oscillates above the surface sample at a given distance. By monitoring the shift of the resonance frequency caused by the attractive forces and maintaining the distance to the surface with a feedback loop, the morphology of the surface can be recorded. It is important to note that the gradient of attractive forces is smaller than the gradient of repulsive forces, meaning that the cantilevers with high resonance frequencies are preferred for a better sensitivity. In tapping mode (Fig.

Conductive AFM

Conductive AFM (C-AFM) is typically done in contact mode and the probe used for this mode of operation has a conductive coating. Fig. 2.10 shows a schematic of the C-AFM set-up. When the probe is in contact with the sample surface, a dc voltage can be applied, then, the current flowing between the conductive probe and the sample is recorded. The electrodes are the probe itself and the metallic sample holder. Note that even if the silver paste is in contact with the perovskite layer, the probe is far enough to avoid any eventual lateral current. 17 The current response signal is recorded simultaneously with the topography of the sample. An external amplifier module (TUNA module from Bruker) is necessary to detect the current with high sensitivity. In this mode, it is also possible to record I-V spectroscopy on specific areas however, the signal to noise ratio use to be high.

Nonetheless, since the current signal from the imaging is reliable enough and constant throughout the measurement, several images done at the same location and different voltages were done in order to reconstruct I-V curves (data presented in Chapter 3).

Kelvin Probe Force Microscopy

Kelvin Probe Force Microscopy (KPFM) is normally used for electrical characterization of interfaces (when using a cross section) and surfaces in the dark or illumination. 19 Is based on the Fermi energy alignment. The Fermi energy, EF, is the energy level where the probability to find an electron with 50 % probability. The energy level occupation by the electrons in the CB and VB is described by the Fermi-Dirac distribution. The difference between the vacuum energy level and the EF is the so-called work function (WF). When the probe is brought close to the perovskite surface, the EF of the material under test and the probe align creating an electric field between them, as shown in Fig. 2.12. 20 Then, the vacuum levels for the probe and the sample are not the same anymore, creating a potential difference (VCPD). A dc voltage (Vdc) is then applied between the probe and the sample to nullify the electric field created by the potential difference, this dc voltage corresponding to the VCPD. As the AFM probe is sensitive to the electrostatic force induced by the tip-sample interactions, KPFM results in a high sensitive technique for CPD determination.

𝑉 𝐶𝑃𝐷 = 𝜙 𝑡𝑖𝑝 -𝜙 𝑠𝑎𝑚𝑝𝑙𝑒 -𝑒
(2.9) The electrostatic force between the tip and the sample (Felec) in the z direction (vertical direction to the sample surface) is given by

𝐹 𝑒𝑙𝑒𝑐 (𝑧) = - 1 2 ∆𝑉 2 𝑑𝐶(𝑧) 𝑑𝑧 (2.10)
where dC/dz is the capacitance gradient between the probe and the surface and ΔV is the difference between the voltage applied to the probe and VCPD, ∆𝑉 = 𝑉 𝑑𝑐 -𝑉 𝐶𝑃𝐷 + 𝑉 𝑎𝑐 sin (𝜔𝑡) (2.11) being Vac the ac voltage with a drive angular frequency ω. Then, introducing Eq. 2.11

into Eq. 2.10, the electrostatic force dependant on whether the bias is applied to the sample or to the probe is expressed as,

𝐹 𝑒𝑙𝑒𝑐 (𝑧, 𝑡) = - 1 2 [𝑉 𝑑𝑐 -𝑉 𝐶𝑃𝐷 + 𝑉 𝑎𝑐 sin(𝜔𝑡)] 2 𝑑𝐶(𝑧) 𝑑𝑧 (2.12)
which can be divided into three components:

𝐹 𝑑𝑐 = - 𝑑𝐶(𝑧) 𝑑𝑧 [ 1 2
(𝑉 𝑑𝑐 -𝑉 𝐶𝑃𝐷 ) being Eq. 2.13 the force of a constant deflection of the cantilever, Eq. 2.14 the force of an oscillation ω and it is used to measure the magnitude of VCPD. Finally, Eq. 2.15

describes an oscillation twice the applied frequency and it is usually used for capacitance measurements. 22 Concerning the determination of the VCPD, it is necessary to determine the difference between Vdc and VCPD. For this purpose, a lock-in amplifier is used in order to extract an electrical force component at the angular frequency ω. Then, a Vdc equal to VCPD is applied, nullifying the output signal from the lock-in. Therefore, the CPD (and also the work function) can be determined all over the surface.

KPFM can be operated in AM (AM-KPFM) or FM modes (FM-KPFM) as in Fig. 2.13.

In AM mode (Fig. 2.13a), the cantilever is excited by a bias sent to the tip with a driving frequency close to the resonance frequency of the cantilever. The CPD is determined from the Fω, measured from the amplitude of oscillation of the probe. Through a PID feedback loop the probe bias is adjusted so the amplitude of oscillation remains minimized, hence the electrostatic force is nullified. The downside of this method of detection is that it measures the all the forces near the probe. Therefore, interactions of the bulk cone from the probe and the cantilever with the surface may be included in the CPD measurements.

VCPD is an averaged value, hence decreasing the resolution of the measurement. In FM mode (Fig. 2.13b), the cantilever is mechanically shaken at its resonance frequency and, at the same time, an ac bias with low frequency is sent to the tip. Then, due to the electrostatic interaction, sidebands appear at ω0 ± ω (Eq. 2.14) and ω0 ± 2ω (Eq. 2.15).

A phase-locked loop is used to measure the frequency shift induced by the perturbation caused by the electrostatic force gradient. A lock-in amplifier is also used to isolate the low frequency electrostatic component. Here, the interaction is confined to the electrostatic forces coming only from the probe apex, which offers an improved resolution compared to AM-KPFM. The main drawback is that FM-KPFM has a lower sensitivity than AM-KPFM and higher ac voltages are necessary, what can make a perturbation in the tip-sample interaction.

Scanning Microwave Microscopy

In order to measure the capacitive behaviour of perovskites at the nanoscale, Scanning

Microwave Microscope (SMM) was used. SMM is composed by a vector network analyser (VNA) and connected to the AFM head by a coaxial cable. Several techniques have been proposed to perform measurements using SMM. [26][27][28] Since the wavelength of microwaves is in the range of cm and the thickness of the device under test (DUT) is in the range of nm, the SMM probes the interaction in the quasi-static near-field. In this work, the VNA is used as both, a wavelength source and a detection system However, the reflection coefficient is measured by the VNA in terms of S-parameters.

Therefore, the VNA measures the so-called S11 coefficient. The S11 is a complex value, varies with the operating frequency and includes information about the amplitude and phase of the measured wave. To define the S-matrix, we need to take into account the incident waves generated in the port of the VNA (incident wave a) and the output wave (or reflected b wave). Then we define S11 as the response of port 1 due to the incident signal in port 1. This nomenclature is especially useful when using two or more ports in the VNA system. Therefore, Γ (or S11) measures intrinsic properties of the DUT such as impedance, capacitance and conductance. Since the probe creates an electric field, the conductivity and dielectric constant of the material can be in principle extracted.

Microwave systems use a standard impedance, Zref = 50 Ω, as a reference. Γ can be expressed as a function of the local (where waves a and b are measured) impedance as:

𝛤 = 𝑆 11 = 𝑍 𝐷𝑈𝑇 -𝑍 𝑟𝑒𝑓 𝑍 𝐷𝑈𝑇 +𝑍 𝑟𝑒𝑓 (2.17)
With a MOS structure where there is a dielectric layer on top of the semiconductor material, the resistance is probably high enough to consider that the capacitance variations take over. In addition, since the measurements are done at the nanoscale, the capacitance is in the order of aF (C = 𝜀•A/d, where 𝜀 is the dielectric constant, A is the area of measurement and d is the thickness). Therefore, the measured impedance range is in the order of MΩ. There is an important mismatch between the impedance of the VNA and the high impedance of the DUT. This means that the measurements done for ZDUT are far from the Zref, lowering the measurement resolution. 29 To overcome this limitation, a 50 Ω lumped resistance is added to the AFM probe holder. The lumped resistance absorbs most of the incident signal, reducing the reflection coefficient to a 0 value. This system allows for accurate measurements for DUT with high impedance in parallel to the 50 Ω resistor. In addition, the sensitivity is increased with two amplifiers, which increase the magnitude of incident and reflected waves. One of the drawbacks of this SMM system is the limited range of measurable DUT impedances, only high ZDUT impedances can be accurately measured. Another drawback is the dependence on the cable length connecting the resistor with the AFM tip, which limits the range of frequencies that can be used. 30 Since it is located at a given distance from the DUT, the cable length is L = n (λ/2) (where λ is the wavelength and n is an integer), the resistor can be considered to be connected to the sample.

A better approach is to use an adjustable interferometer instead, as in Fig. 2.17. 31 Fig.

2.17 represents the SMM configuration with a VNA (PNA in the scheme) and the AFM represented by the "SMM probe" in the scheme. Here, when the wave is split in two, the reference signal is adjusted in magnitude by the variable attenuator and then combined with the reflected wave b (or aref in the scheme) in a different point of the circuit (Mach-Zehnder interferometer model). At this point, regardless of the impedance magnitude, the resulting wave can be nullified with the attenuator. With the cable length used in our setup, the interference occurs more or less every 98.1 MHz, so the interference can be adjusted in a range of ± 50 MHz. This allows for a wider range of frequency operation compared to the method with a lumped resistance. Finally, the wave is amplified before being collected by the receiver. SMM is also very useful for measuring the local capacitance. The capacitance can be measured in two ways: it can be derived from the S-parameter or it can be directly measured with an additional differential capacitance (dC/dV) module added to the interferometer. For the first method, several calibration procedures have been reported but here we used two calibration methods: the calibration procedure described by Keysight Technologies and an interferometric model developed at the IEMN. 32 The second method to measure capacitance consists of a direct measurement of the dC/dV.

EFM calibration

For a given frequency, the measured S11 can be expressed in terms of error parameters (e00, e01 and e11) and related to the S parameter of the probe-sample system S11,a as:

𝑆 11 = 𝑒 00 + 𝑒 01 𝑆 11,𝑎 1-𝑒 11 𝑆 11,𝑎 (2.18)
where the error values include spatial losses, the stray capacitance, the phase shift and the attenuation produced by the couplers and splitters. While some error parameters, such as the one produced by the attenuation, depend mainly on the SMM system and remain constant through measurements, other errors such as the one produced by the stray capacitance, are made in-situ, that is, they are associated to the different tip-sample distance in each experiment. Generally, the error parameters are determined from three known reference samples but, since the stray capacitance will be disturbed when changing the samples, Gramse et al. implemented a method that allows determining the error values from the DUT. 32 They used an approach-retract curve and recorded the S11 though the entire tip-sample distance. For dielectric samples, the impedance is only capacitive and can be measured at low frequencies (~ 1.5 kHz) by electric force microscopy (EFM). As described in section 2.2.6, the second harmonic is used for capacitive measurements thus it is recorded here. Although in EFM the second harmonic could be recorded, it constitutes a complex problem. 33,34 EFM involves van der Waals interactions between the probe and the sample, which are affected by any contaminants on the surface and/or the atmospheric conditions. It also modulates the capacitance with the vibration mode of the cantilever giving increased contribution from the moving end of the probe. This may limit the accuracy of the calibration. However, in contrast to EFM, SMM provides additional capabilities for capacitance measurements. In particular, since it operates in contact mode, an additional DC voltage can be added to evaluate the dependence of the capacitance with that voltage and deduce material properties such as doping type and eventually doping level.

Following Eq. 2.10, the integration of the electrostatic force vs the applied voltage curve results in the capacitance C(z). Next, from the approach-retract curve combining EFM and S11, the error values can be determined. The capacitance has different contributions:

tip-sample, cone-sample and probe plateau-sample. Near the sample, at a distance in the range of the tip radius, due to the local electrical interactions, the tip-sample contribution is large and follows a non-linear trend. In contrast, far from the sample the capacitance variation becomes small compared to the overall capacitance and can be linearized, as observed in Fig. 2.19. 35

Interferometric calibration

To convert the S11 into complex impedance Z, an interferometric calibration is under development. First, a frequency for the measurement is set near a given frequency interference. Then, the procedure consists on recording the interference of the S parameter over two different frequency ranges. One range will span over several interferences, while the second frequency range will focus only on the interference near the frequency we set for the measurement.

Considering that the sinusoidal wave ai created in the VNA source can be written as

𝑎 𝑖 = 𝑎 0 𝑒 𝑖𝜔𝑡 (2.19)
where a0 is a complex number composed by its amplitude |a0| and phase φa0 at the initial point, x = 0. Then, at a given point of the system, the wave is split in two where one is directed to the DUT and the other one is used as a reference. Therefore, now it is necessary to take into account the distance (x) to the source at a given time (t):

𝑎(𝑥, 𝑡) = 𝑎 0 𝑒 -𝑖(𝑘𝑥-𝜔𝑡) (2.20)
where k is the wave vector (k=ω/νg). Once the incident wave reaches the DUT and is reflected, it is defined as:

𝑎(𝑥, 𝑡) = 𝑒 10 𝛤𝑎 0 𝑒 -𝑖(𝑘𝑥-𝜔𝑡) (2.21)
where e10 includes the spatial losses, the phase shift and the attenuation produced by the couplers and splitters. The term Γ represents the reflection coefficient after interacting with the DUT. The other wave is directed to the attenuator and is defined as:

𝑏(𝑥, 𝑡) = 𝑒 20 𝑎 0 𝑒 -𝑖(𝑘𝑥-𝜔𝑡) (2.22)
where, again, e20 includes the spatial losses, the phase shift and the attenuation. The wave coupling is made in a given point, called x1 for the reflected wave and x2 for the reference wave. The amplitude of both waves is added following the expression:

𝑎(𝑥, 𝑡) + 𝑏(𝑥, 𝑡) = 𝑒 10 𝛤𝑎 0 𝑒 -𝑖(𝑘𝑥-𝜔𝑡) + 𝑒 20 𝑎 0 𝑒 -𝑖(𝑘𝑥-𝜔𝑡) (2.23) and then amplified by a complex gain, G, as:

𝑎 𝑟 (𝑥, 𝑡) + 𝑏 𝑟 (𝑥, 𝑡) = 𝐺𝑎 0 (𝑒 10 𝛤𝑒 -𝑖𝑘𝑥 1 + 𝑒 20 𝑒 -𝑖𝑘𝑥 2 )𝑒 𝑖𝜔𝑡 (2.24)

The combined signal is detected by the VNA, which corresponds to the S parameter:

𝑆 = 𝐺(𝑒 10 𝛤𝑒 -𝑖𝑘𝑥 1 + 𝑒 20 𝑒 -𝑖𝑘𝑥 2 ) (2.25)
which corresponds to the complex amplitude from the combined wave. Since the length of the 2 different cables for the incident wave and the reference wave is not the same, there is a phase shift between them. This allows to nullify the S parameter using the attenuator, which can adjust e20, resulting in a virtual Γ0 used as a reference. When the signal is nullified, which occurs at some given angular frequency of interest 𝜔 0 corresponding to wavelength 𝑘 0 , 𝑒 20 𝑒 -𝑖𝑘 0 𝑥 2 = -𝑒 10 𝑒 -𝑖𝑘 0 𝑥 1 𝛤 0 , and then:

𝑆 𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑 = 𝐺𝑒 10 𝑒 -𝑖𝑘 0 𝑥 1 (𝛤𝑒 -𝑖(𝑘-𝑘 0 )𝑥 1 -𝛤 0 𝑒 -𝑖(𝑘-𝑘 0 )𝑥 2 ) (2.26) However, even in the case where the interference is not perfectly nullified, a virtual reflection coefficient Γ0 nullifies the combined signal:

𝛤 0 = - 𝑒 20 𝑒 10 𝑒 -𝑖𝑘 0 (𝑥 2 -𝑥 1 ) (2.27)
Eq. 2.26 can be expressed as:

𝑆 𝛤 (𝑘) = 𝑆 𝛤 0 (𝑘) + 𝐺𝑒 10 𝑒 -𝑖𝑘 0 𝑥 1 (𝛤 -𝛤 0 )𝑒 -𝑖(𝑘-𝑘 0 )𝑥 1 (2.28)

Eq. 2.28 can be rewritten with the frequency difference (f -f0) and time difference (t1 -t2) instead of the wave vector difference (k -k0) and spatial position difference (x1 -x2), respectively. Using also 𝐴 = 2𝐺𝑒 10 𝑒 -𝑖𝑘 0 𝑥 1 = |𝐴|𝑒 𝑖𝜑 𝐴 , the first term becomes:

𝑆 𝛤 0 (𝑘) = |𝐴|𝛤 0 sin[𝜋(𝑓 -𝑓 0 )(𝑡 2 -𝑡 1 )]𝑒 𝑖[ 𝜋 2 -2𝜋(𝑓-𝑓 0 ) 𝑡 1 +𝑡 2 2 +𝜑 𝐴 ] (2.29)
where Γ0 is a reflexion coefficient corresponding to an admittance that can be virtually described by a capacitance (C0) and a conductance (G0). The second term is rewritten as:

𝑆 𝛤 (𝑘) -𝑆 𝛤 0 (𝑘) = |𝐴| 𝛤-𝛤 0 2 𝑒 𝑖[-2𝜋(𝑓-𝑓 0 )𝑡 1 +𝜑 𝐴 ] ] (2.30)
Due to the small probing area in SMM measurements, the admittance is usually small, therefore, Γ ≈ Γ0 ≈ 1. This means that usually the term

𝛤-𝛤 0 2 𝑒 𝑖[-2𝜋(𝑓-𝑓 0 )𝑡 1 +𝜑 𝐴 ] is much lower than sin[𝜋(𝑓 -𝑓 0 )(𝑡 1 -𝑡 2 )].
When the experimental data is measured over a frequency range such that (𝑓 -𝑓 0 )(𝑡 1 -𝑡 2 ) ≅ 2, they can be fitted with Eq. (2.29), which makes it possible to determine t1 -t2 and t1 + t2 from the periodicity of the amplitude and phase over several periods. Also, fitting the experimental interferences allows to obtain the values of amplitude and phase (|A| and φA) so that the fitting is best near the interference near the frequency of interest (f0), see Fig. 2.20a and b. However, due to the non-linear dependence of the electrical properties with the frequency, the frequencies far from f0 do not fit well. Nonetheless, the curve follows well the experimental data in a range of around With t1 and t2 estimated, now we use a frequency range much shorter than before, as in Fig. 2.20c and d so as to fine tune the value of A near the interference. We track the amplitude and phase variation over typically hundred MHz around the f0. If we consider f0 ≈ f, Eq. (2.29) can be written as:

𝑆 𝛤 0 (𝑘) = -|𝐴|𝛤 0 𝜋(𝑓 -𝑓 0 )(𝑡 1 -𝑡 2 ) 𝑒 𝑖[ 𝜋 2 -2𝜋(𝑓-𝑓 0 ) 𝑡 1 +𝑡 2 2 +𝜑 𝐴 ] (2.31)
where |A| and 𝜑 𝐴 can be found by fitting 𝑆 𝛤 0 (𝑘)) vs frequency. Note that in Fig. 2.20, the fit is done using Eq. (2.28) for better validation of the model. This allows in particular to fine tune the dependence of 𝑆 𝛤 (𝑘) in the vicinity of f0 (10 kHz range) by taking into account the admittance (Y0) that makes the interference not fully abrupt (phase shifting by 180° without continuity due to the change of sign of (f-f0)).

So far, |A| was determined from the wide frequency measurement and from the shorter frequency measurement, both giving similar values yet slightly shifted. This problem can be improved by enhancing the microwave components. Both t1 and t2 have been also estimated. Therefore, we can proceed to measure a given sample in order to determine its admittance.

The determination of the sample's admittance (and capacitance) can be done if we take into account that the frequency set to make the measurement (f1) may be shifted from f0 by a small quantity (a few tens of kHz for instance), so that:

𝑆 𝑘 1 (𝛤) ≅ 𝐵 𝛤-𝛤 1 2 ≅ 𝐵𝑍 𝐶 (𝑌 1 -𝑌) (2.34)
where 𝛤 1 = 𝛤 0 𝑒 -𝑖(𝑘 1 -𝑘 0 )(𝑥 2 -𝑥 1 ) , 𝐵 = 𝐴𝑒 -𝑖(𝑘 1 -𝑘 0 )𝑥 1 and Y1 is the virtual admittance that gives 𝑆 𝑘 1 (𝛤 1 ) = 0 at f1. 𝑌 1 -𝑌 is here assumed to be << 1/ZC. With the calibrated parameters, the admittance can be deduced.

Capacitance-Voltage Measurements

The C-V curve can be obtained by measuring the S11 for a range of dc biases. Fig. 2.21a

shows an ideal C-V curve for a p-type semiconductor. This curve can be understood by taking the analogy with a metal-oxide-semiconductor (MOS) capacitor. In this work, the metal is the AFM probe (Rocky Mountain Nanotechnology xxPtyyyA), the oxide is Al2O3 and the semiconductor material is the perovskite under study. The capacitance associated to each region can be determined using a parallel plate approximation where the MOS structure can be seen as a group of capacitors connected as in Fig. 2.21b. 38 In the flat band region, the energy bands (conduction and valence bands) in the semiconductor are flat. This is reached between the accumulation and depletion behaviour in Fig. 2.21a. It corresponds to the case where the charge in the MOS capacitance is null.

This situation is reached by applying a voltage (flat band voltage) that compensates the work function difference between the metal probe and the semiconductor. However, if there are charges in the oxide, the flat band voltage shifts to positive or negative values, depending on the sign of the charges. 39 Additional charges may also be present at the semiconductor interface that may be charged or discharged depending on the applied voltage. They compete with the semiconductor depletion, which results in a dilatation of the voltage scale in Fig. 2.21a. 38 In the accumulation region, the majority charges (holes in p-type) are driven towards the anode (tip side) creating an accumulation at the oxide-semiconductor interface. This phenomenon occurs when applying negative voltages to the probe. The accumulation region can be further understood looking at the scheme in Fig. 2.21a, the total capacitance (Ctot) is mainly the oxide capacitance (Cox) and it is independent on the doping type and carrier concentration of the semiconductor:

𝐶 𝑡𝑜𝑡 = 𝐶 𝑜𝑥 = 𝜀 𝑜𝑥 𝜀 0 𝐴 𝑡 𝑜𝑥 (2.35)
where 𝜀ox is the electric permittivity of the oxide, 𝜀0 is the dielectric permittivity in the vacuum, A is the surface area of the probe contact and tox is the oxide thickness.

When applying positive voltages to the probe, the majority carriers are driven towards the cathode (semiconductor side), this is known as depletion region and it induces a space region with a characteristic capacitance (Cdepl), this is represented in Fig. 2.21b. The total capacitance will depend on both the Cox and the Cdepl:

𝐶 𝑡𝑜𝑡 = 𝐶 𝑜𝑥 𝐶 𝑑𝑒𝑝𝑙 𝐶 𝑜𝑥 + 𝐶 𝑑𝑒𝑝𝑙 (2.36)
where

𝐶 𝑑𝑒𝑝𝑙 = 𝜀 𝑠𝑒𝑚 𝜀 0 √ 2 𝜀 𝑠𝑒𝑚 𝜀 0 𝜙 𝑠𝑒𝑚 𝑞 𝑁 (2.36)
being ϕsem the interface potential of the semiconductor, q the elementary charge and N is the acceptor dopant concentration of the semiconductor. If the positive voltage is further increased, then the energy bands will bend further downwards eventually crossing the fermi energy level, creating what is called "inversion". At this point, the capacitance depends on minority carriers. Depending on the ability of minority charges to follow the applied frequency will result in either a higher capacitance (normally for low frequencies)

or lower capacitance (for high frequencies). At low frequencies, the system is in thermal equilibrium and minority carriers can respond, therefore only Cox is taken into account:

𝐶 𝑡𝑜𝑡 = 𝐶 𝑜𝑥 (2.37)
When the frequencies are high, minority carriers cannot follow so they do not contribute to the total capacitance and it decreases to a level below the depletion capacitance.

Therefore, Ctot will be a combination of Cox and Cdepl at its maximum width:

𝐶 𝑡𝑜𝑡 = 𝐶 𝑜𝑥 𝐶 𝑖𝑛𝑣 𝐶 𝑜𝑥 + 𝐶 𝑖𝑛𝑣 (2.38)
where

𝐶 𝑖𝑛𝑣 = 𝐴 𝜀 𝑠𝑒𝑚 𝜀 0 √ 2 𝜀 𝑠𝑒𝑚 𝜀 0 2 𝜙 𝑠𝑒𝑚 𝑞 𝑁 (2.39)
However, the stray capacitance induced by the system is usually in the order of picofarads (pF), which is much larger than the sample's capacitance (in the order of attofarads, aF).

Therefore, a good technique to get rid of parasitic capacitance is the differential capacitance measurement.

Differential capacitance measurement

Another way to measure the capacitance is through a direct measurement of the differential capacitance (dC/dV) at different dc biases, similar as in scanning capacitance of the SMM system using the dC/dV module in blue. The scheme represents the SMM configuration as in Fig. 2.17 plus the addition of the dC/dV module in blue. The dC/dV module can be mainly considered as an amplifier to be able to detect the small variations. The measurement of dC/dV has been shown to be especially useful in the semiconductor industry where non-linear behaviour at low frequency (RF) signal results in an impedance variation. Even if the impedance variation is typically low, the difference between the incident VNA signal and the low frequency cause a signal magnification. The relationship between the capacitance difference (ΔC) and the impedance difference (ΔZ) is:

∆𝑍 = 1 2 𝜋 𝑗 𝑓 ∆𝐶 (2.40)
where f is the microwave operation frequency. Therefore, the sensitivity is a prori increased compared to direct low frequency measurement and will further increase for higher frequencies. 40 The shift in dC/dV will be more significant around the depletion range, where the capacitance-voltage (C-V) curve will show the largest slope possible,

Introduction

Organic-inorganic hybrid perovskites have been capturing much interest in the field of photovoltaics (PV) as the power conversion efficiencies (PCEs) of corresponding devices have been strongly improved, from 3.8% in 2009 to more than 23% recently. 1,2 This prodigious increase in just a decade is the result of intense research and fundamental understanding of crucial intrinsic properties of the perovskites; in particular, their electronic and charge transport properties feature remarkably long diffusion lengths and carrier lifetimes. [3][4][5][6][7][8] However the device stability and the mechanisms responsible for this low stability are still under debate. 9,10 In addition to the bulk properties, the surface properties, which depend on the crystallographic facets and the possible presence of defects, are also of prime importance as their impact on the electronic and transport properties at grain boundaries or other interfaces is significant, hence affecting the performances of the devices. [11][12][13][14][15][16][17][18] Several methods are proposed for reducing these detrimental effects for photovoltaic performances, involving, among others, tuning of the perovskite composition and interfacial engineering. [19][20][21] For instance, reducing the recombination at the interface can be achieved by introducing a chemical linker between the perovskite and the electron transport layer. 22 Recombination sites were also reduced after passivation of grain boundaries using either an excess of methylammonium iodide (MAI) during the synthesis or a post-fabrication surface treatment. 23,24 Yet, to face such a major issue, the electronic properties of the surface of perovskite thin films ought to be further investigated where variations can be observed, i.e. at the nanoscale. In this regard, Scanning Probe Microscopy (SPM) provides several characterization methods to locally map the electronic properties of materials with nanometer resolution. 25 between electrons and holes. 27 Harwell et al. exposed the MAPbI3 perovskite repeatedly to light and found that the Fermi level eventually returns with time to its initial value with slow decay, this effect being attributed to trapped or slow charges within the device. 33 Stecker et al. used SPM to identify vacancy-assisted transport and, combined with density functional theory (DFT) calculations, they predicted an increase of the work function when increasing the number of vacancies in MAPbBr3. 34 Gallet et al. found electronic heterogeneities at the MAPbI3 perovskite surface and attributed them to differences in surface terminations. 35 From these works, it has been mostly agreed that among the surface features ruling the electronic properties of the perovskite, defects or traps are those responsible for the electronic heterogeneities observed at the local scale. Traps can be classified as shallow or deep, depending on the energy position in the gap. Calculations on the activation energies of all the possible point defects as determined from those studies revealed that usually shallow trap states are more likely to form than deep traps, and they can be an explanation for the long electron-hole diffusion lengths observed in these materials. 36 Yet, Du et al. identified iodine interstitials, with specifically low activation energies that could generate significant deep trap states. 37 And all these studies are also pointing out how the type and density of traps at the surface of the perovskite are depending on the growth conditions of the materials along with the anisotropy of the surface properties. In spite of different types of defect evidenced at the surface, their corresponding impact onto the film stability are still under debate. To enhance perovskite stability, it is therefore crucial to understand how these defects specifically affect the surface properties and how they interact with water molecules, especially at the early stage of degradation.

In this chapter, frequency modulated KPFM (FM-KPFM) and Conductive Atomic Force Microscopy (C-AFM), are used to evidence the presence and the electronic contribution of trap states at the surface of methylammonium lead iodide (CH3NH3PbI3 or MAPbI3) perovskite thin films. FM-KPFM is an appropriate method to simultaneously probe topographic and surface potential variations without cross-talking. 25,29 It measures the contact potential difference (CPD, i.e. the work function difference) between the metallic probe and the sample under test, hence providing meaningful insights towards device performances and stability, see Chapter 2. 38 C-AFM was shown to be pertinent for highresolution electrical characterization of perovskites. 39 As a current measurement method, C-AFM can be sensitive to contributions from surface properties, such as defects, as they are expected to impact the charge transport mechanism via carrier injection. 23,39,40 Combining FM-KPFM and C-AFM allows us to locally correlate carrier injection to work function (WF) in view of understanding the electronic properties of the sample surface.

Characterization of the CH3NH3PbI3 Surface

To gain insight into the perovskite films, we deposited MAPbI3 on top of a glass/ITO/PEDOT:PSS substrate, using PbAc2 as a precursor and following the fabrication procedure described in section 2.1 for MAPI-PbAc2. 41,42 Freshly-prepared MAPI-PbAc2 thin films were then analyzed by XRD and XPS to determine their chemical composition and crystallinity, respectively, in order to establish reference profiles prior to any degradation process, and to confirm standards reported in literature for pristine MAPbI3. As shown by the XPS spectrum of Fig. 3.1a, the major atomic species present at the surface of the MAPI-PbAc2 are, as expected, iodine, nitrogen, carbon and lead. The additional presence of indium atoms could be attributed to the presence of pinholes in the PEDOT:PSS + MAPbI3 stack. 43 The corresponding XRD profile, displayed in Fig. 3.1b, indicates a dominating presence of (110) and equivalent (220) crystalline planes corroborating the high crystal orientation expected from this fabrication route. 44 Moreover, the characteristic peak of PbI2 expected at ~ 12.6° (see asterisk symbol in Fig. 3.1b), and usually related to perovskite phase segregation, is absent, allowing these XRD and XPS data to serve as reference profiles for pristine MAPI-PbAc2. 45 The corresponding morphology of the films is composed of grains of varying size (170 ± 130 nm in diameter), as determined with non-contact AFM (Fig. 3.1c) operated in an inert environment. The topography of the films reports a typical RMS roughness around 34 nm for 4×4 µm² images. was obtained from the surface potential data after tip calibration, see section 2.2.6).

Fig. 3.2b

shows spatial heterogeneity of the WF at the film surface, as highlighted by the complex WF distribution curve in Fig. 3.2d. Note that the surface properties are studied at the MAPbI3/N2 interface, however, there may be substantial changes on the interface properties when the contact with MAPbI3 is made with other gas (like oxygen) or films used for device applications. The substrate on which the perovskite is grown and/or operating conditions (i.e. temperature or pressure) may also affect the surface properties.

The WF distribution can be deconvoluted into two dominating populations, which have an average difference of ~ 0.1 eV. A high WF domain (HWF) and a low WF domain (LWF) are highlighted with white and red circles in Fig. 3.2b, respectively. However, if we take consider the WF values only at the center of the grains, those two domains show up to 0.2 eV WF difference; 4.80 ± 0.02 eV and 5.00 ± 0.02 eV for LWF and HWF respectively. This calculation is made by masking the HWF and LWF domains in the raw image, as in Fig. 3.3a and b. Masks were created by selecting the regions which possess values above or below a given threshold value, thus, HWF and LWF domains appear highlighted respectively. Results will then be discussed as if there were two clearly distinguished populations, although since the deconvoluted profiles are quite large and overlap, there might be intermediate populations lying between and outside the dominating contributions. The schematic representation of LWF and HWF on the KPFM configuration can be seen in Fig. 3.4.

The measured WF values lie in a range between 4.70 and 5.10 eV. The reported electron affinity and ionization energy for MAPbI3 are most commonly around 3.9 and 5.4 eV, respectively, although there is some scatter in the literature. [46][47][48][49] This would indicate that despite the heterogeneous WF at the surface of the film, the Fermi level remains located between the valence band edge and mid-gap, what means that the whole surface of the MAPI-PbAc2 thin film is exhibiting p-type semiconducting behavior, hence with holes as majority carriers. 46 Notice that the topographic profile does not influence nor correlate with the CPD signal (see Fig. Because FM-KPFM is probing surface properties, the observed correlation between FM-KPFM and C-AFM images strongly supports that the C-AFM response is also related to the perovskite surface properties. In other words, the C-AFM current is ruled by perovskite surface properties (i.e. the local work function) at the tip-sample contact.

So far, we observed different surface domains associated to given local work functions and, at the same time, we found them to be related to different current densities. Surface heterogeneity is further investigated by means of averaged C-AFM images taken at different biases, where I-V profiles are built for the different domains. Fig. 3.6a displays the C-AFM current variations with bias at these two locations, exhibiting an exponential law between 1 and 2 V in both domains, consistently with a specific injection/extraction dominating mechanism. In this bias range, the I-V profiles are characteristic to that of a rectifying diode, suggesting for both locations a Schottky-like contact at the tip-sample contact with an associated built-in potential (Vbi) in the films. MAPI-PbAc2 exhibiting ptype behavior, this diode-like contact is therefore biased in forward polarity upon positive dc sample bias.

For a standard Schottky contact, the charge transport between the probe and the sample in forward bias is conditioned by the thermionic-emission theory, which analytical expression is:

𝐽 = 𝐽 0 𝑒 𝑞𝑉 𝑛𝑘𝑇 (1 -𝑒 -𝑞𝑉 𝑘𝑇 ) (3.1)
Where J is the current density, J0 is the saturation current density obtained in reverse bias, q is the charge of an electron, V is the voltage applied to the sample, n is the ideality factor of the diode, k is the Boltzmann constant and T is the temperature; 50 and with J0 ∝ 𝑒 -𝑞ϕ 𝑏 𝑘𝑇

where ϕb is the effective barrier of the Schottky contact. The ideality factors, determined from the slope in the semilog profiles of Fig. 3.6a, are similar. The difference in current between the two types of domains in the semilog I-V curve can therefore be mainly attributed to different effective barriers (Δϕb). Fig. 3.6b and 3.6c illustrate this casescenario, where holes must overcome a higher barrier in LWF domains before they can be collected by the metallic probe.

We also note that a divergent field might be expected for the probe-sample contact geometry. This effect would be present in case of an Ohmic contact, however, we observed rather a Schottky behavior. 51 Here we assume that the current is limited to the electronic properties of the contact, which will mainly affect to the magnitude of the measured current (~ nA). The current is limited to the injection mechanism and, considering the semiconducting properties of the perovskite, thermionic emission was assumed as the dominant mechanism in forward bias. The dotted arrows illustrate the fitting of the bandgap at the X-axis intercept.

The variations of WF and ϕb between HWF and LWF domains are also consistent with the p-type semiconducting character of MAPbI3, lower (higher) WF inducing higher (lower) Vbi. Generally, the WF and free carrier density (p) are conditioned by the sole crystallinity of the film and play a primary role onto charge collection, determining the injection voltages. 53,54 In this work, however, charge injection/extraction mechanisms are dominating the current response up to much higher injection voltages, indicating significant amount and impact of surface states, hence possible interfacial issues once incorporated in devices. 55,56 This is likely to induce a Fermi level pinning located at the energy level of the traps (4.8 and 5.0 eV as previously shown). In C-AFM, the alignment of the Fermi level is taking place between the probe and the sample in both the LWF and HWF domains with their respective trap energy levels at 0 V. This Fermi level pinning has an impact on the probe-sample contact and can be described with an additional electric field across an interfacial layer as depicted in Fig. 3.6b and c.

The existence of domains exhibiting different surface states despite similar stoichiometry can be attributed to, at least, two kinds of different perturbing states; they could either be the result of different terminations or the result of the growth procedure. Recent DFT calculations gave insight into the effect of the possible different MAPbI3 terminations. 57 Pure methylammonium(MAI)-terminated (001) surfaces are around 1 eV higher in energy than pure PbI2-terminated (001) surfaces, both independently affecting the charge injection properties. 58,59 In addition, MAI-terminated surfaces are expected to have a higher bandgap than PbI2-terminated surfaces. 59 However, given the fact that the two types of domains observed here exhibit similar VB, it is quite unlikely that they would correspond to two different surface terminations. We are therefore led to hypothesize that surface structural defects induced by the growth conditions would cause the observed local electronic heterogeneity. However, at this point we cannot completely discriminate between these two possibilities. The presence of defects will be further examined in Chapter 5.

Alternative CH3NH3PbI3 Perovskite

So far, the MAPI-PbAc2 perovskite showed a p-type behavior. In fact, MAPbI3 thin films can be self-doped depending on the ratio between the precursors used in the synthesis process. 60 High (low) concentrations of PbI2 in the precursor solution makes the final perovskite film to be more n-type (p-type), respectively. Moreover, PEDOT:PSS has a better affinity to PbI2 than to MAI, so even higher amounts of MAI are necessary to generate a p-type MAPbI3 perovskite on the PEDOT:PSS surface. 61 Previously, we used MAPI-PbAc2, the synthesis of which results in a p-type material. To prepare a n-type perovskite, we used PbI2 as the precursor, in a 1:1 (PbI2:MAI) ratio. It is referred as MAPI-PbI2 (see section 2.1) and it corresponds to one of the most typical synthesis routes used in MAPbI3 perovskite solar cells. 62,63 The MAPI-PbI2 perovskite surface in this case shows a surface with small grains of ~ 150 nm, Fig. 3.8a. This perovskite surface was also characterized by KPFM and C-AFM. The anisotropy in WF and current images is only apparent, that is, the highlighted domains in The bandgap energy is ~ 1.56 eV (Fig. 3.9a and b), in line with the bandgap reported in literature and, with the previously mentioned energy levels, it implies an n-type behavior for this perovskite film. 3 The approach previously used to reconstruct an I-V curve was also tempted for the MAPI-PbI2 film in order to confirm the n-type behavior. This technique requires to make several C-AFM scans of the same area at increasing bias voltages after each scan. However, the described procedure is only apparently simple. A negative current was found when scanning at 0 V (note that the measurements are carried out in dark conditions), what means that there is a small parasitic photocurrent, most probably due to the AFM laser beam (680 nm wavelength) used to control the AFM probe motion. In addition, for each C-AFM scan (at different bias voltages) there is a small portion of the image scanned at 0 V, this allows us to check the "base" current prior to apply any voltage. Nonetheless, the negative current for this 0 V tests is increasing after each scan. This odd behavior translates into an increasing photocurrent contribution after each scan and a misleading I-V curve. However, based on (i) our experimental protocol and (ii) what is regularly and consistently reported in literature, we assume these MAPI-PbI2 films to exhibit n-type semiconducting properties. Altogether, the MAPI-PbI2 surface exhibits homogeneous electronic properties (excluding the fact that there are unreacted PbI2 species) over the surface compared to the MAPI-PbAc2 surface. In MAPI-PbI2, deep traps were also found, since the average surface WF was ~ 4.4 eV. In this case, since the WF difference is almost negligible, we could assume a certain specific acceptor point defect to be the major defect present on the surface. The discussion about the possible defect candidates will be further discussed in Chapter 5.

Conclusions

In contrast to previous literature reports, we have demonstrated that even if MAPI-PbAc2 perovskite films show a very high crystalline structure seen by XRD diffractograms, they can show regions with up to 0.2 eV WF difference. We discriminated between two different regions, HWF and LWF, where LWF regions have a surface WF closer to the mid-gap, hindering carrier injection/extraction at the surface, hence the charge transport in the film. We proposed either surface terminations or surface point defects to be the cause of these electronic local variations. Since the bandgap is expected to remain unaltered through the perovskite surface, it is unlikely that surface defects are causing such a WF perturbation. We then suggest a heterogeneous distribution of point defects over the MAPI-PbAc2 surface. In contrast, when using lead iodide as a precursor (MAPI-PbI2), the WF distribution over the surface appears to be much more homogeneous.

Interestingly, even though the WF is more homogeneous in this perovskite film, the average surface WF indicates the presence of deep traps too. The determination of the surface traps will be further discussed in Chapter 5 devoted to the evolution of the surface under water contamination.

Introduction

The remaining uncertainty about the chemical properties at the perovskite surface can be a potential factor that limits the approaches when trying to improve the perovskite-based devices. Therefore, it is essential to fully understand what processes are taking place at the perovskite surface to, later, understand the behavior at the interface between the perovskite and a given transport layer.

One of the limitations when measuring surface properties is the size of the measurement.

Measuring electronic local properties in the order of few nanometers presents a challenge and it becomes harder when several measurements of different nature are required in the same nanometer-sized spot. While it exists a number of SPM-based techniques to test small features in perovskite devices, as described in Chapter 2, studies reporting local capacitance properties are still very limited. Impedance spectroscopy measurements have been shown to be a powerful method to determine the capacitance and the processes taking place at different parts of the device, yet it remains a macroscopic measurement. 1,2 For photovoltaic applications, the main interest in capacitive measurements is to be able to know more about conductivity, diffusion lengths and/or carrier lifetimes.

In order to understand the local capacitive effects, scanning microwave microscopy (SMM) could be a useful technique. SMM can detect small variations in the reflected microwave signal and it has been used in the microelectronics industry, for instance, to distinguish different doping levels. 3 However, while it has been proven to be a reliable technique to detect local capacitive features, meaningful studies with perovskite-based devices are lacking. 4 In this chapter, preliminary SMM measurements have been carried out in order to study the local capacitive variations on both routes previously described in MAPbI3 seen in following the EFM calibration procedure on section 2.2.7 -EFM calibration for S11 measurements and, as a result, a linear trend was approximated from Fig. 4.2b. Notice that after the probe contacts the sample, the capacitance keeps increasing, this might be due to geometrical changes in the probe tip or due to the bending of the cantilever. In order to minimize the topographical contributions on the SMM measurements, the following linear regression was systematically used for this purpose: 4 In Fig. 4.3a and b, the |S11| can be observed before and after crosstalk correction, respectively. The correction allowed to decrease the signal from grain boundaries, which use to be problematic points for the probe-surface interaction and result in small artifacts.

Even after crosstalk correction, it becomes evident that there is a contrast over the perovskite surface. In the |S11| histogram, Making work function and |S11| images at the same spot could lead us to distinguish between the different populations. Such a study could be done by different ways, one of them consists to mark a nanometric spot with an AFM tip so when the WF map is measured, the spot can be found later when moving the sample to the SMM system.

Nonetheless, the camera in the SMM system does not provide sharp enough images,

Calibration for Capacitance Measurements

The S11 signal can be converted to capacitance following the EFM calibration procedure described in section 2.2.7 -EFM calibration. EFM calibration has been published and it is extensively used in literature, however, it is necessary to have the capability to make EFM and S11 measurements at the same time in order to get a calibrated sample. 6 This is actually a limitation for this work. S11 measurements were first performed in the Keysight SMM system where EFM measurements could be done through the AFM integrated capabilities, however, another SMM system was built by the IEMN with the added advantage of being operated under vacuum conditions. In this later system, the EFM measurements were not integrated just yet. Therefore, an alternative calibration technique was needed.

Calibration is a fundamental aspect generally in all kind of measurements, however, especially here it is a limiting factor. The most straightforward method to calibrate the measurements would be to measure the S11 signal in another well-known sample of similar characteristics, and it is a usual method usually performed with samples such as silicon or other materials. 7,8 Nonetheless, with other materials such as perovskites, it is complicated to find a reference sample with similar characteristics. In addition, the fact of using several samples implies that the probe-sample distance will be different in each case, therefore, sometimes it is not the most reliable way to calibrate the S11 measurements. IEMN developed a calibration method, so-called interferometric calibration, in which the S11 signal is recorded in a range of frequencies once the probe is in contact with the sample and prior to make any measurement, see section 2.2.7 -Interferometric calibration for further details. This technique does not need any additional measurement mode.

Therefore, the interferometric calibration was tested together with the EFM calibration just by measuring an approach-retract curve on a MAPbI3 perovskite sample. In Fig. 4.4

the results show negative capacitance values for the interferometric model. However, the results for the interferometric model had to be adjusted by adding an additional capacitance of around 1 pF in order to show a correct trend, that is, a higher capacitance close to the surface. This additional capacitance is rather high compared to what it is expected to be measured on the sample (~ fF), therefore, it should be due to some parasitic effect due to the instrumentation, probably caused in the SMM head. It could be that the propagation of the high frequencies between the coaxial cable and the sample play an important role. Such complex propagation schemes may be difficult to solve and would require electromagnetic 3D modelling (COMSOL Multiphysics for instance). Even though, regardless of the uncertainty of the capacitance values calculated using the interferometric calibration, it can provide a qualitative analysis which can be helpful to distinguish the local capacitive behaviour without needing to use an additional reference sample nor any additional AFM mode. Therefore, this will be the used calibration technique for the capacitance measurements in the next sections.

Local Capacitance

To further study the origin of the |S11| contrast seen in Unfortunately, the resolution of the measurement and the small difference between domains prevents us to draw any reliable conclusion with dC/dV measurements.

Nonetheless, from Chapter 3 we know that MAPI-PbAc2 shows heterogeneous electronic properties due to different kind of defects. Therefore, these local capacitance images serve as preliminary results leading us to think that defects present on the surface are creating a capacitance difference, probably due to a difference in their defect densities.

Since we could not combine KPFM and SMM at the same location, a C-V curve could ultimately help us to distinguish between n or p-type behavior in the different perovskite surfaces: MAPI-PbAc2 and MAPI-PbI2. To perform this study the SMM system was used under vacuum conditions so no insulating layer was necessary, see Fig. depending on the working conditions while growing the perovskite, the contrast may change significantly. Also, the resolution for this study was decreased down to 64 pixels 2 , so the measurement may not detect eventual local variations as much as in Fig. 4.5

(resolution of 128 pixels 2 ). The resolution was decreased in order to decrease the time needed to make multiple S11 maps. Nonetheless, the purpose of this study was to test the capabilities of the SMM system and to track any capacitance change that the perovskite may show upon voltage variation.

Then, the capacitance was averaged for the MAPI-PbAc2 maps at different applied bias voltages, Fig. 4.7a. Also, the C-V measurements were done twice as to be sure about the reproducibility, being both measurements very similar. A remarkable divergence between forward and reverse measurements can be easily noticed, a priori indicating a hysteresis effect. Given the hysteretic evidences in MAPbI3 perovskites (see section 1.3.3), it lead us to think that the capacitance is mostly coming from the perovskite layer. Nonetheless, the current is very small especially for LWF regions, therefore, a hysteresis effect may be harder to detect than in SMM.

Conclusions

In this chapter, we first showed how to extract topographical artefacts from the SMM measurements by using a linear regression. The local S11 contrast has been also studied and compared to the work function contrast. While we found similarities in both SMM and KPFM contrasts we would need a deeper analysis to conclude they are arising from the same phenomena. The interferometric calibration has been tested and compared to the EFM calibration, providing a similar behaviour but with the values shifted. This is something that could be further investigated by making 3D models of the wave propagation from the end of the probe to the sample. Still, using the interferometric calibration for qualitative comparison, C-V curves have been made, showing a higher hysteresis on MAPI-PbI2 samples compared to MAPI-PbAc2 samples. After comparing them to I-V curves in the same range of voltages, an unusual behaviour at positive voltages in MAPI-PbI2 has been detected, thus, leading us to think that ion motion, from species such as Ii, is causing the high hysteresis effects.

As a summary it has been shown that SMM is a promising tool for capacitance characterization at the nanoscale. The results show: (1) contrasted capacitance images with several domains associated to a capacitance variation, and (2) the detection of a bigger MAPI-PbI2 perovskite hysteresis compared to the MAPI-PbAc2 perovskite.

Introduction

As previously described in Chapter 1, perovskite solar cells are well known for their impressive optoelectronic properties. Their ease of fabrication also makes them to be in an economical advantage over other technologies such as silicon-based solar cells. Their latest improvements have brought them to the same performance level as other technologies such as thin-film photovoltaics. [1][2][3][4] However, the main reason holding them back to be competitive in the solar cell market is their lack of long term stability. 5 In section 1.3.5 the stability in perovskite was widely described, in fact, stability is usually addressed macroscopically either for the whole solar cell device or for the perovskite layer itself. 6 Measurements such as photoluminescence, current-voltage (I-V) curves, quantum efficiency and many other related macroscopic measurements have been proven

to be useful for the understanding of the device behavior under different conditions.

Nevertheless, there still lacks a link between the observed macroscopic stability and the stability at the local scale. Lately, some groups are trying to push their perovskite research into that direction. For instance, Garret et al. used KPFM under illumination and found a reversible ion migration taking place at the MAPbI3 perovskite surface, making the Voc to vary upon illumination. 7 Such properties at the nanoscale may have affect the overall device performance and/or stability. Nonetheless, most of the perovskite stability studies are typically performed in ambient condition or in a humid environment. In this chapter we track the local stability with KPFM and, at the same time, we introduce a more precise method allowing for a controlled low water exposure level. For this study, MAPbI3 was prepared by the two different synthesis routes previously described in section 2.1 and characterized in Chapter 3. The SPM measurements, carried out for both sets of samples, exhibit different surface electronic properties and stability upon exposure to water.

Complementary X-ray photoelectron spectroscopy (XPS) and X-ray diffraction (XRD)

analysis correlate the observations with the variations in the chemical composition and the crystalline orientation planes, respectively, for the two types of MAPbI3 perovskite.

Consistent pictures can be drawn between chemical structure and composition, surface states and electronic properties and water stability for each of the two differently synthesized perovskites. The purpose of this work was to study the evolution of the work function and current of a given location in the perovskite surface after water exposure. However, such study implies that the sample has to be moved from the AFM stage to the vacuum chamber for water exposure. Then, the sample needs to be placed again in the AFM stage for the measurement of the work function (by KPFM) and current (by C-AFM). Measurements with KPFM and C-AFM require different AFM probes, what implies that the AFM probe holder would necessarily be moved. Therefore, a method to identify the same micrometric spot was needed. The approach we followed is presented as follows:

Sample Preparation for Tests

First, the sample is marked with a cross using a tweezer, which makes it easy to spot with the AFM camera. Still, the cross is quite big compared to the probe size, so we need to make some smaller marks. Next, the topography is scanned near the cross mark in contact mode, taking care of not scratching the sample. Once a suitable location is found, the probe is lifted by decreasing the setpoint voltage so that the AFM feels like if it was in contact. Actually, this makes the AFM to scan while lifted. At this moment, we make the AFM to scan in just 1 horizontal line (we tend to use 512 lines for a whole image)

repeatedly. Then we engage again but this time the setpoint voltage is set to be very high so that the sample surface is significantly scratched by the probe. After some time, the probe is lifted and the same process is made for a parallel line with around 7 µm of separation. This results in two parallel lines (and a destroyed probe) easily noticeable with the AFM camera, as in Fig. 5.1. Next, it is just a matter of looking for the small lines with the AFM camera and, once engaged, to place the probe in the area between the lines by detecting the borders of it when tracking the topography.

Impact of Small Amounts of Water on the CH3NH3PbI3 Perovskite Surface

Therefore, the WF was measured taking care of always probing the same location These results indicate that for the MAPI-PbAc2 films, very small exposure to water is enough to strongly affect the electronic properties by creating deep trap states, hence to suppress charge transport. In this scenario, in a pristine p-type perovskite, the presence of surface states makes the surface to be less p-type, as previously shown in Fig. 3.4, the possible point defects to be considered are donor sites lying above the Fermi level. These defects can be induced by I vacancies (VI), substitution of Pb and MA cations at I sites (PbI and MAI, respectively), substitution of Pb ions at MA sites (PbMA) and MA and Pb interstitials (MAi and Pbi, respectively). 13 Among these defects, we consider MAi, and VI to potentially dominate due to their lower formation energies. 14,15 We also investigated the effect of water exposure on the perovskite prepared from the PbI2 precursor (MAPI-PbI2). When exposing the MAPI-PbI2 surface to pulsed water vapor cycles, the WF remains unchanged up to 200 cycles (Fig. be dehydrated back to MAPbI3. 16 For comparison, in our tests we exposed the samples to much lower amounts of water and, in addition, the vacuum chamber was heated at ~ 55 ºC, which would dehydrate the intermediate back to MAPbI3. After prolonged exposure to water vapor, the perovskite can transform into a dihydrated compound ((MA)4PbI6•2H2O) that cannot be dehydrated back. 16 Here, MAPI-PbAc2 and MAPI-PbI2

perovskites cannot be measured anymore by SPM techniques after 120 and 200 cycles respectively. The dihydrated perovskite compound may be creating electrostatic interactions through the coordinated H2O and the AFM probe. The work function of MAPI-PbI2 is located closer to the mid-gap than to the conduction band, the trap states observed in MAPI-PbI2 surface are then expected to be acceptor ones inducing deep states. These surface acceptor defects seem to be less reactive upon water exposure, therefore we assume n-type surfaces such as MAPI-PbI2 are more stable at the early-stage of water-induced degradation.

Tracking the Structural Evolution at the Early Stage of Water Degradation

Both pristine MAPbI3 materials were also characterized by XPS, and similar measurements were carried out after exposure to 40 and 100 water vapor cycles, corresponding to 9 x 10 8 L and 2.2 x 10 9 L. However, when transferring the samples from the glovebox to the XPS, it is inevitable to expose the samples to ambient conditions for a short period of time. The relative humidity of the atmospheric air in our lab is around 50%. The saturated vapor pressure of water at 21 ºC is 18.650 torr. 17 If we take into account that the time needed to transfer the samples from the glovebox to the XPS chamber is around 1 minute: Although this is a considerable amount of water, it is still below the amounts of water vapor injected in the vacuum chamber during 40 and 100 water vapor cycles. When looking at the amount of oxygen present in both MAPbI3 films (Fig. 5.2e and f), it seems that the amount of water is greatly affected by the water vapor injected in the vacuum chamber rather than by the ambient exposure. Therefore, we considered that this fast exposure to ambient conditions has little effect for these measurements. To elucidate the specific interactions of water with MAPI-PbAc2 and MAPI-PbI2, we recorded high-resolution XPS core-level spectra. The C 1s core level spectrum of MAPI-PbAc2 shows two clear peaks at 285.1 eV and 286.8 eV, for C-C and C-N groups respectively, Fig. 5.7a. 20 Even though the peak at 285 eV has been typically attributed to C-C bonding and it is used for XPS calibration, it is most probably a bond created from small amount of hydrocarbon contamination. 21,22 However, this peak has been also assigned to CH3I, created when MAI vapor is in contact with PbI2. 23,24 After 40 and 100 cycles of water, the C 1s photoemission peaks and the relative carbon concentration remain unchanged, as observed in Fig. 5.7c, 6e, and 4e. Since the nitrogen content was shown to decrease after water exposure along with the increase of the oxygen content, the peak at 286.8 eV for MAPI-PbAc2 can be attributed to a combination of contributions from C-N groups and C-O groups. The evolution of the oxygen and nitrogen content is consistent with the occurrence of a nucleophilic substitution reaction by which the oxygen atom of a water molecule binds to the carbon atom of methylammonium, expelling one ammonia molecule that is pumped away. Although this reaction is kinetically not favorable, we think that it is enhanced by the vacuum chamber heat (~ 55ºC). This chemical reaction is expected to deeply affect the electronic properties of the materials.

Note however that the increase in the oxygen content is much stronger than the decrease in the nitrogen content (Fig. 5.2e), which suggests that part of the water molecules simply adsorb on MAPI-PbAc2 (and could further modify its electronic properties via electrostatic interactions).

The C 1s core level spectrum of the MAPI-PbI2 shows two peaks at ~ 285 eV and ~ 287 eV, similar to the MAPI-PbAc2 peaks. However, not only the C 1s peaks and the relative concentration of carbon remain unaffected, but also the concentration of oxygen and nitrogen remains unchanged upon water exposure, as shown in Fig. 5.7b, d, f and Fig. 5.2f. This observation confirms that MAPI-PbI2 is less reactive with water. 25 The Pb 4f emission shows two distinctive peaks at 138. the peak for metallic Pb can be also found (~ 142 eV), Fig. 5.8f. The presence of metallic lead in the perovskite surface has been previously attributed to the photolysis of the unreacted PbX2 materials (X being an halide) induced by the X-rays. 26 In particular, in this work, we identified unreacted PbI2, see the peak with an asterisk symbol in Fig. 5.4c.

Therefore, it is not surprising to find metallic Pb as previously reported. 27 Comparing again our results with the work from Ralaiarisoa et al., they detected a reduction of the metallic Pb content, suggesting that this reduction introduces gap states that pin the Fermi level, however, in this study the lead content remained constant even after exposing the samples to water vapor. 12 The constant behavior was also measured for the I 3d content, with emission peaks at 619.

Conclusions

In this chapter, a relationship between electronic structures at the MAPbI3 surface and local stability upon low water exposure has been proposed. In MAPI-PbAc2, the WF decreases after water exposure. This effect was shown to correspond to a drift of the Fermi energy towards mid-gap, indicating the creation of deeper trap states turning the perovskite into a semi-insulator, consistently with the collapse of the current signal. The XPS measurements indicate that the methylammonium groups are likely to react with water molecules releasing NH3 and creating C-O bonds. Also, a higher stability against moisture was observed for MAPI-PbI2. This improved stability was attributed to the different nature of defects, acceptors compared to donors in MAPI-PbI2 and MAPI-PbAc2 respectively. By using a vacuum chamber, we had a precise control on the exposure to very small amounts of water, which allowed us to study the early degradation mechanisms in perovskites. We observed that strong interaction with water molecules can occur very quickly in p-type perovskites like MAPI-PbAc2. This rapid interaction with water alters the electronic and transport properties even prior to any morphological or stoichiometric changes.

Conclusions

In this thesis, the main objective is to track the surface properties of perovskites at the nanoscale by a characterization process including scanning probe microscopies and surface analysis. These surface properties play an important role for obtaining efficient and stable perovskite solar cells. However, since the properties of the surface are significantly related to the synthesis route and the environment conditions, it is important to fine-tune each synthesis step. Therefore, in this work, we combined several characterization techniques in order to give some insights about the resulting surface properties with respect to two different synthesis routes. This chapter summarizes the main conclusions of this work, and then provides some suggestions for future research towards perovskite solar cells with a better understanding of the surface properties.

First, the CH3NH3PbI3 perovskite was studied at the nanoscale using AFM techniques such as KPFM and C-AFM. The analysis focused on two different synthesis routes for CH3NH3PbI3 so-called MAPI-PbAc2 and MAPI-PbI2. These synthesis routes lead not only to a different perovskite bulk character but also to different surface properties. The MAPI-PbAc2 perovskite route leads to a p-type perovskite while MAPI-PbI2 results in an n-type perovskite. KPFM maps revealed an heterogeneous MAPI-PbAc2 surface in terms of work function (WF). We approximated the WF map to be composed mainly of two domains. Among the possible origins of the WF contrast, we considered two: either surface terminations or surface point defects. However, the effective barriers measured for the two domains, together with their respective WF allowed us to determine a similar valence band (VB) for each domain. Given the fact that PbI-and MAI-terminated surfaces differ in their energy levels and also their bandgap by ~ 1 eV, we suggest point defects to be responsible for the WF difference. We also measured a lower WF at the surface compared to the bulk of the perovskite, meaning that among the possible point defects, donor sites are probably the ones responsible for the anisotropy. MAPI-PbI2 perovskite was also analysed by KPFM and, putting aside the fact that there were unreacted PbI2 species, no major WF contrast was observed over the surface. Here, with an average WF of 4.4 eV and an n-type perovskite, the possible traps to be considered are acceptor ones.

We further characterized the perovskite electronic surface properties using the SMM.

After the removal of the topography crosstalk signal on the S11 signal, we were able to detect a significant contrast in S11 over the surface. To study it further, the S11 was converted to capacitance variation by using a calibration method based on an interferometer. Capacitance variation also showed heterogeneous properties over the MAPI-PbAc2 surface, similar to what was observed by KPFM and C-AFM in Chapter 3.

Next, the average capacitance in both perovskite surfaces (MAPI-PbAc2 and MAPI-PbI2) was tested upon bias voltage. The result was a big hysteresis effect especially for MAPI-PbI2 surfaces, probably due to a higher ion motion induced by the characteristic Ii surface defects that can be found in n-type perovskites such as MAPI-PbI2.

Finally, with the main purpose to test the stability of these two perovskite films, they were exposed to water in a perfectly controlled system. The exposure of a very small amount of water on MAPI-PbAc2 resulted in both domains (before mentioned) changing their WF towards lower values. The fact that both domains shift their WF in an even trend further excludes the hypothesis of surface terminations as the origin behind the WF difference as they are expected to interact differently to water. Another important observation is that no current could be detected anymore from the first water exposure, meaning that the perovskite turns into a semi-insulator. After repeatedly exposing the MAPI-PbAc2 surface to increasing amounts of water, the WF value reached a plateau around mid-gap. This effect suggests that donor surface defects are interacting with water. XPS measurements confirmed the presence of oxygen on the MAPI-PbAc2 surface. Same as for MAPI-PbAc2, the MAPI-PbI2 surface was also exposed to water vapour. However, this time the electronically uniform MAPI-PbI2 surface kept its average WF for much higher exposures than MAPI-PbAc2. This may be an indication of acceptor sites being less reactive than donor sites upon water interactions.

Outlook

As it is the case with research studies, ideas are developed in a given time, which is usually not enough to bring them fully to completion. Here there are some outlooks for future pathways of research that could also further improve the work done in this thesis.

So far, we have determined how synthesis routes, and more specifically, how the p and n semiconducting character can affect the stability of perovskite films. This study leaves room for a wider local stability study in which the much sought after device stability

Figure 1 . 1 .

 11 Figure 1.1. Installed power generation capacity by source in the Stated Policies Scenario, 2000-2040. 2

Figure 1 . 3 .

 13 Figure 1.3. Global total solar photovoltaic market scenario until 2020 predicted by SolarPower Europe.4

Figure 1 . 4 .

 14 Figure 1.4. Evolution of the solar cell performance for the different photovoltaic technologies.18

A

  sites are generally Cesium (Cs) or organic cations, B sites are either Pb or Sn and are surrounded by anion X within a [BX6/2] octahedral framework. X sites are usually halides.

Figure 1 . 5 .

 15 Figure 1.5. ABX3 Perovskite crystal structure.

Figure 1 . 6 .

 16 Figure 1.6. (a) UV-Vis absorption spectra of FAPbIyBr3-y perovskites. (b) Solar spectrum recorded on the International Space Station.38,39

53 Figure 1 . 7 .

 5317 Figure 1.7. Perovskite solar cells with (a) n-i-p and (b) p-i-n structures.

Figure 1 . 8 .

 18 Figure 1.8. Scheme of the basic mechanism in a perovskite solar cell: (1) light absorption and exciton generation, (2) exciton migration and (3) charge carrier transfer and extraction to the electrodes.

Figure 1 . 9 .

 19 Figure 1.9. J-V curves of a solar cell under dark (blue curve) and illumination (red curve) conditions with the main points labelled. The grey rectangle represents the fill factor.

Figure 1 . 10 .

 110 Figure 1.10. Forward bias to short circuit (FB-SC) and short circuit to forward bias (SC-FB) current-voltage curves measured for the same halide perovskite-based solar cell.

Fig. 1 . 11 .

 111 However, the VB does have a strong anticoupling bonding between Pb s and I p states, raising the VB energy above the I p orbital level. The electronic states are then affected when substituting the halide, for instance, the bandgap increases from 1.5 eV in MAPbI3 to 2.3 eV in MAPbBr3 due to the smaller size of Br ions, which also originates in a different electronic energy distribution. 101 Defects induce additional states which could be positioned between the CB and VB, thus, affecting the Voc and device performance. Leijtens et al. investigated the high Voc observed in perovskite solar cells and observed that trapped electrons possess long lifetimes, only recombining with free holes over the course of many microseconds.

Figure 1 . 11 .

 111 Figure 1.11. Schematics of the VB and CB in CH3NH3PbI3 perovskites. 103

Figure 1 . 12 .

 112 Figure 1.12. Calculated transition energy levels of donor-like and acceptor-like point defects in MAPbI3. The possible defect charges are noted in parentheses at the band level position. Also in parentheses are the formation energies for the neutral defects.103

Figure 1 . 13 .

 113 Figure 1.13. Schematic model showing the capture of excited electron and formation superoxide. Due to favourable energy offset, some electrons can be accepted easily in the case of m-TiO2 (left) and due to the high energy difference between the perovskite and the mesoporous Al2O3, electrons cannot be injected into perovskite instead are captured by molecular oxygen resulting in the formation of superoxide (right).110

Figure 1 . 14 .

 114 Figure 1.14. In-situ HR-XPS spectra of both pristine and thermally treated MAPbI3. The graphic shows the iodine to lead (I/Pb) and nitrogen to lead (N/Pb) ratio in both pristine and thermally treated samples. 124

Figure 1 . 15 .

 115 Figure 1.15. XRD analysis of pristine and degraded CH3NH3PbI3. Black squares, red circles, and green stars shows PbI2, pristine CH3NH3PbI3 and hydrated CH3NH3PbI3 respectively. 127

Fig. 1 .

 1 16a. 141 Fullerenes help to passivate the perovskite surface and reduce the effects of deep traps. 150 Most common materials used as HTLs are poly(3,4-ethylenedioxythiophene) polystyrene sulfonate (PEDOT:PSS), 2,2',7,7'-Tetrakis[N,N-di(4-methoxyphenyl)amino]-9,9'spirobifluorene (Spiro-OMeTAD) and poly(triarylamine) (PTAA), see Fig. 1.16a, c, b

1 .Figure 1 . 16 .

 1116 Figure 1.16. Chemical structure of: (a) PCBM, (b) Spito-OMeTAD, (c) PEDOT:PSS, (d) TBP, (e) LiTFSI and (f) PTAA.

3 . 5 :

 35 ITO/PEDOT:PSS/Perovskite. However, in some parts of Chapter 4, an insulating layer of Al2O3 is added after the perovskite deposition in order to avoid degradation due to the characterization technique, which are made in atmospheric conditions First, transparent and patterned indium-tin-oxide (ITO) substrates (Naranjo substrates, ref. NS1463, 95 nm, ~ 20 Ω.sq) are cleaned thoroughly. The surface is wiped down with soap to remove any microscopic contaminants such as organic compounds, dirt and dust. Then the substrates are ultra-sonicated with detergent, deionized water (15 MΩ/cm 2 ), acetone and isopropanol for 10 minutes each. They are then dried under a stream of nitrogen gas. Finally, the samples are O2-plasma cleaned for 5 min. After the cleaning process, a thin layer of PEDOT:PSS is deposited. PEDOT:PSS (Heraeus Clevios P Al 4083) is spin-coated at 5000 rpm for 60 s and baked at 110 °C for 10 min. The final thickness of the PEDOT:PSS layer is around 40 nm. Then, the samples are transferred into an N2-filled glovebox in order to deposit the MAPbI3 perovskite, see Fig. 2.1. At this point, two different synthetic routes were carried out.

Figure 2 . 1 .

 21 Figure 2.1. Glovebox used for (a) perovskite synthesis and (b) deposition.

7 Figure 2 . 2 .

 722 Figure 2.2. Four perovskite deposition processes: (a) one-step precursor deposition, (b) two-step sequential deposition, (c) dual-source vapour deposition and (d) vapourassisted solution process.8

Figure 2 . 3 .

 23 Figure 2.3. Schematic conductivity type conversions in perovskite films by excess CH3NH3I (left side) or excess PbI2 (right side). Possible point defects in perovskite films caused by composition variation were illustrated correspondingly. The elemental defects

  diffractogram. The diffracted peaks can be associated with different planes in an atom structure and they are named after the (hkl) Miller indices. In the XRD set-up, the sample under test is exposed to a beam of x-rays with a given incident angle θ as in Fig.2.4. The diffracted light is received by the X-ray detector in a 2θ angle. Diffraction is used to measure the distance between the atom planes.

Figure 2 . 4 .

 24 Figure 2.4. Schematic representation of the Bragg's equation.

  However, for p, d and f peaks, two peaks are observed (spin orbital splitting). The XPS technique was used to understand the chemical composition of the perovskite sample surface (in Chapters 3 and 5). XPS measurements were carried out with a PHI 5300ESCA PerkinElmer spectrometer and the recorded spectra were calibrated with the C 1s peak (285.1 eV).

Figure 2 . 5 .

 25 Figure 2.5. Photoemission process in XPS. Black dots represent electrons and the bars represent energy levels in a given atom.

2 . 6 )

 26 by the equation:𝛼ℎ𝛾 = 𝐴(ℎ𝜈 -𝐸 𝐺 ) 1/2(2.8)UV-visible absorption spectra was taken with a Lambda 650 S PerkinElmer spectrometer.

Figure 2 . 6 .

 26 Figure 2.6. UV-Vis spectra and Tauc plot for the MAPI-PbAc2 perovskite with an optical bandgap estimated around 1.59 eV.

Fig. 2 . 7 ,

 27 and it is divided into different components. The AFM probe is located at the end of a cantilever, which has a given spring constant (0.1 to 20 N/m typically). The spring constant affects directly the force applied to the sample; soft samples could be damaged with probes with high spring constant. The motion of the probe is controlled by a piezoelectric actuator and a feedback loop maintaining the amplitude modulation to a user-defined value. The piezoelectric actuator vibrates at the resonance frequency of the cantilever under an applied ac bias voltage, it will produce the cantilever oscillation. The light of a laser is projected on the back of the cantilever, which is reflected to a four-quadrant photodiode, Fig.2.7.

Figure 2 . 7 .

 27 Figure 2.7. Scheme of an AFM set-up.

2 .

 2 8c), the cantilever is kept at a certain distance from the sample and, it oscillates in a given amplitude touching the sample. The amplitude is controlled by an ac signal applied to the piezo at, or close to, its resonance frequency. The amplitude of oscillation depends on the interaction between the probe and the surface. It is kept constant though a feedback loop, which allows for a non-destructive scanning of the surface under test. The measurements are based on electrostatic forces, establishing a direct relationship between the force of the probe and the sample. The relationship between force and probe-sample distance can be seen in Fig.2.8d.

Figure 2 . 8 .

 28 Figure 2.8. Atomic force microscopy modes: (a) contact, (b) non-contact or frequency modulated AFM and (c) Tapping mode or amplitude modulated AFM. (d) Typical forcedistance curve in AFM.

Figure 2 . 9 .

 29 Figure 2.9. Bruker Multimode AFM inside a N2-filled glovebox.

Figure 2 . 10 . 2 . 11 .

 210211 Figure 2.10. Schematic of the C-AFM set up. CAS stands for current amplifier setup.

Figure 2 . 11 .

 211 Figure 2.11. Force curve in Peakforce mode. Red letters represent the important information points on the curves.

Figure 2 . 12 .

 212 Figure 2.12. Electronic energy levels of the sample and AFM tip for three cases: (a) tip and sample are separated by distance d with no electrical contact, (b) tip and sample are in electrical contact, and (c) external bias (Vdc) is applied between tip and sample to nullify the CPD and, therefore, the tip-sample electrical force. Ev is the vacuum energy level. Efs and Eft are Fermi energy levels of the sample and tip, respectively.21

Figure 2 . 13 .Fig. 2 . 14 .

 213214 Figure 2.13. Schematic diagram of KPFM system showing (a) AM and (b) FM mode.23

Figure 2 . 14 .

 214 Figure 2.14. Connections between the AFM Multimode, the Signal Access Module, the Nanoscope controller, the Oscillation Contol 1 (OC1) and the Nanoscope computer.

Figure 2 . 15 .

 215 Figure 2.15. Variation of the frequency shift with the probe-sample distance.25

  . The VNA normally contains a radio frequency source generator (RF), a splitter, a detector and a digital processor. It measures the amplitude and phase magnitudes of small reflected electrical signals. The SMM used in this work operates from 2 to 18 GHz (AT-SMM from Keysight Technologies). Fig. 2.16 shows the diagram of a VNA. The VNA splits the generated wave in two, one wave is used as the reference signal and the other is sent to the DUT, referred to as the incident wave. The reflected wave is directed to the measurement channel. At the source side, an attenuator gives the ability to change the output power depending on the response of the DUT.

Figure 2 . 16 .

 216 Figure 2.16. Scheme of a two port VNA.

Figure 2 . 17 .Fig. 2 .

 2172 Figure 2.17. SMM set-up integrating an adjustable microwave interferometer based ona March-Zender configuration.31

Figure 2 . 18 .

 218 Figure 2.18. SMM systems: (a) IEMN-made and (b) Keysight technologies.

Figure 2 . 20 .

 220 Figure 2.20. Comparison between the experimental interference curve around 18.21 GHz and the modelled curve using Eq. 2.28 for (a) (b) a wide range of frequencies and (c) (d) a shorter range of frequencies. The used sample was composed of ITO/PEDOT:PSS/MAPI-PbAc2.

Figure 2 . 21 .

 221 Figure 2.21. (a) Ideal C-V curve of n-type semiconductor. (b) Metal-Oxide-Semiconductor (MOS) scheme of the SMM set-up with the associated capacitances.

  microscopy (SCM). As a comparison, in SCM two signals are simultaneously analysed: one in the GHz frequency range and the other one in the kHz range. Instead, the SMM can discriminate ac signals since only the reflected wave with the same frequency as the incident wave is detected by the VNA. In fact, SMM can measure dC/dV by applying an ac voltage with a low frequency (in the kHz range) and a microwave signal coming from the VNA (2 to 18 GHz) sent through an additional external module. The differential capacitance is then measured with respect to a fixed dc bias. Fig. 2.22 shows the diagram

Figure 2 . 22 .

 222 Figure 2.22. Diagram of the VNA (black area) and dC/dV module (blue area)components.40

Fig. 2 . 23 .

 223 Fig. 2.23. Under these conditions, high dC/dV values indicate low carrier concentration or low doping (red curve in Fig. 2.23) while low dC/dV values correspond to high carrier concentration or high doping. The signal is then analysed with a lock-in amplifier. The set up used in this work allows to use signals in a range of 50 kHz to 500 kHz.

Figure 2 . 23 .

 223 Figure 2.23. Ideal dC/dV vs voltage curves derived from C-V curves.

  -31 Kim et al. recently used Kelvin Probe Force Meicroscopy (KPFM) and bias-dependent atomic force microscopy (AFM) to show that the properties of the (112) surface of a perovskite crystal can be affected upon external bias while the (100) surface remains unchanged. 32 KPFM has also been used by many other groups to study the surface potential of perovskite thin films: Bergmann et al. used cross-section KPFM on a MAPbI3 based solar cell to show the unbalanced charge-carrier extraction

Figure 3 . 1 .

 31 Figure 3.1. Characterization of freshly-prepared MAPI-PbAc2 deposits: (a) XPS survey spectrum, (b) X-ray diffractogram (the asterisk symbol indicates the angle corresponding to the reflection expected for PbI2) and (c) typical non-contact AFM height image (cropped from a 4x4 µm image to see more easily the grains).

Fig. 3 .

 3 Fig. 3.2a and b show the topography and WF, respectively, as measured by FM-KPFM for a MAPI-PbAc2 film deposited on glass/ITO/PEDOT:PSS substrate (the WF image

3 .Fig. 3 .Fig. 3 .

 333 Fig. 3.2e does not show 2 clear peaks (as for the work function, Fig. 3.2d), the correlation between current and WF is further evidenced by the averaged scatter plot shown in Fig. 3.5c. The current variation with the topography was meanwhile found to be inside the error bar when compared to the height variations (see Fig. 3.5b), indicating little if no correlation also between these two signals.

Figure 3 . 2 .

 32 Figure 3.2. SPM images of MAPI-PbAc2 deposited on glass/ITO/PEDOT:PSS substrate showing the (a) topography, (b) work function (measured with FM-KPFM), (c) current image (measured with C-AFM in Peak-Force tunneling TUNA mode, taking the peak current with an applied dc sample bias of 1.7 V), (d) work function histogram (with the color grading used in (c)) and (e) current histogram. Note that the FM-KPFM and C-

Figure 3 . 3 .

 33 Figure 3.3. Masks of the (a) HWF and (b) LWF regions used to determine the average WF in these domains.

Figure 3 . 4 .

 34 Figure 3.4. Schematic representation of the LWF and HWF domains in KPFM.

Figure 3 . 5 .

 35 Figure 3.5. Scatter plots of (a) the height vs. CPD, (b) the height vs. current and (c) the work function vs. current for MAPI-PbAc2 perovskite.

Figure 3 . 6 .

 36 Figure 3.6. (a) Semi-log plot of the bias-dependence of the C-AFM current (measured in contact mode) for the LWF and HWF types of MAPbI3 domains. (b) and (c) Scheme of the C-AFM band structure of the probe-surface contact for the two types of domains, LWF and HWF at 0 V. The surface states (represented by the red dot) impact the position of the Fermi level at the MAPbI3 surface. There is an interface layer between LWF/HWF and the metallic probe which allows for vacuum level adjustment.

Figure 3 . 7 .

 37 Figure 3.7. (a) UV-Vis absorption spectra and (b) Tauc plot of the MAPI-PbAc2 sample.

Fig. 3 .

 3 Fig. 3.8a, b and c are associated to the small peak at zero current observed in the C-AFM current distribution, Fig. 3.8e, which corresponds to insulating PbI2 domains. The presence of PbI2 domains are confirmed by the small peak at 12.6º in the XRD

Figure 3 . 8 .

 38 Figure 3.8. (a) Height, (b) work function and (c) current (applied bias: 1.5 V) maps of a MAPI-PbI2 layer deposited on top of PEDOT:PSS. Distribution of the (d) work function map and (e) the current map. (f) XRD diffractogram of the MAPI-PbI2 sample. Red circles show the same spot.

Figure 3 . 9 .

 39 Figure 3.9. (a) UV-Vis absorption spectra and (b) Tauc plot of the MAPI-PbI2 pristine and water exposed samples. The dotted arrows illustrate the fitting of the bandgap at the X-axis intercept.

Chapter 3 .

 3 Two different systems were used, on one side the SMM 5600LS by Keysight Technologies, in the other side, the IEMN-made SMM allowing us to work under vacuum conditions and to avoid moisture degradation.

|𝑆 11 |

 11 = 𝛼 + 𝛽 1 * 𝑧(𝑥, 𝑦) + 𝛽 2 * 𝛻 2 𝑧(𝑥, 𝑦) (4.1)where z (x,y) is the topography height, 𝛻 2 z (x,y) is the Laplacian of the topography and α, β1 and β2 are determined from the regression method. The Laplacian operator was introduced to correct the local curvature that is typically found in AFM scans. Fig.4.1dshows the same |S11| vs topography height after applying the linear regression.

Figure 4 . 2 .

 42 Figure 4.2. (a) Schematics of the capacitance contributions between the probe and the sample, (b) Capacitance dependence on the probe-sample distance, the dotted line represents when the probe contacts the sample.

Fig. 4 .

 4 3c, we could distinguish 2 populations at 0.223 dB and 0.229 dB. This contrast is interestingly similar to the contrast observed for the same MAPI-PbAc2 surfaces in Chapter 3, plotted in Fig. 4.3d. In Chapter 3, the discussion was around the possibility to have different surface defects, leading to a work function difference.

  making it very difficult to find the nanometric spot. Another way to perform such study would be to measure simultaneously the KPFM and SMM signals. Nonetheless, each technique is performed with different AFM probes; the resonant frequency for the KPFM probes used in this work is ~ 75 kHz while for SMM probes it is ~ 9 kHz. Therefore, either upgrading the camera system or making an AFM probe that could work with both modes would really represent a big improvement for SMM measurements.

Figure 4 . 3 .

 43 Figure 4.3. (a) Raw and (b) corrected |S11| signals from the MAPI-PbAc2 surface. (c) |S11| histogram and (d) work function histogram for the MAPI-PbAc2 perovskite surface (two different samples).

Figure 4 . 4 .

 44 Figure 4.4. Capacitance measurements calibrated with the EFM and interferometric models.

Fig. 4 . 3 ,

 43 the capacitance values were determined following the before-mentioned interferometric calibration model. The resulting capacitance image, Fig. 4.5b, shows a heterogeneous capacitance over the MAPI-PbAc2 surface. The capacitance heterogeneity may be due to a variation in the surface perovskite properties such as: carrier density, carrier type or defects.

Figure 4 . 5 .

 45 Figure 4.5. (a) Topography, (b) local capacitance image determined from the corrected S11 signal and (c) differential capacitance image on the MAPI-PbAc2 surface.

Fig. 4 .

 4 Fig. 4.5a and b, respectively. While the topography shows the typical grain size for the MAPI-PbAc2 synthesis route (Fig 4.6a), the capacitance (Fig. 4.6b) seems to be less contrasted compared to the previous MAPI-PbAc2 layer shown in Fig. 4.5b. In fact,

Figure 4 . 6 .

 46 Figure 4.6. (a) Topography and (b) derived capacitance from the S11 signal for a MAPI-PbAc2 surface.

Figure 4 . 7 .

 47 Figure 4.7. (a) Capacitance and (c) current vs voltage in two different MAPI-PbAc2 perovskite samples.

Figure 4 . 9 .

 49 Figure 4.9. (a) Capacitance and (b) current vs voltage in another MAPI-PbAc2 perovskite surface.

Figure 5 . 1 .

 51 Figure 5.1. Marks done with an AFM probe for spot tracking.

(Fig. 5 .

 5 2a, b). Upon exposure to water, the WF of the MAPI-PbAc2 surface tends to decrease for the two types of domains, HWF and LWF (further described in Chapter 3), towards a common value, as shown in Fig. 5.2c. After 9•10 8 L (40 cycles) of water exposure, HWF and LWF domains can no longer be distinguished and, upon further water exposure, the WF tends to stabilize at ~ 4.6 eV. Note that at such low exposure, the topographic profile remains unchanged (Fig. 5.2a). The final WF value corresponds to the Fermi level being in the vicinity of mid-gap. For a non-perfect crystal, such intrinsic behavior (having the Fermi level at mid-gap) can only be attributed to semi-insulating properties, where the energy bands are depopulated from their carriers by deep traps, inducing a weak conduction. This makes the perovskite to behave as an insulator.

Figure 5 . 2 .Figure 3 . 4 .

 5234 Figure 5.2. Work function images of the MAPI-PbAc2 (a) pristine film and (b) after exposure to 100 cycles of water vapor. Graph (c) shows the WF evolution of MAPI-PbAc2 while graph (d) shows the WF evolution for MAPI-PbI2 upon water exposure. Graph (e) and (f) represent the atomic concentrations measured by XPS for MAPI-PbAc2 and MAPI-PbI2, respectively. Merged domains (blue dots in (c) and (d)) stand for the situations where domains can no longer be distinguished so only the average WF value of the sample is considered.

5 .

 5 2d). The C-AFM current also remain constant up to 200 cycles (see Fig. 5.3). This is in strong contrast with the MAPI-PbAc2 material, for which no current could be measured after the first 10 water vapor cycles (corresponding to 2.25•10 8 L). Interestingly, Leguy et al. also investigated the initial steps of MAPbI3 degradation by exposing the films to 35% relative humidity at 21 ºC and found that MAPbI3 forms hydrated intermediates (MAPbI3•H2O), which can

Figure 5 . 3 .

 53 Figure 5.3. Current response average measured in C-AFM at 1.3V for MAPI-PbI2 upon exposure to water cycles. The decrease of the current beyond 200 cycles is thought to be due to formation of a layer of condensed water on the surface which results in MAPbI3 perovskite segregation.

Fig. 5 .

 5 Fig.5.2e shows the variations of the atomic concentration of the elements forming the MAPI-PbAc2 films. Upon exposure to water, a significant increase of the oxygen content at the surface can be observed and, simultaneously, a decrease in nitrogen concentration.In contrast, the nitrogen concentration barely changes in MAPI-PbI2 upon water exposure while no significant amount of oxygen is observed (Fig.5.2f). It is tempting to relate these behaviors to the WF shift towards the mid-gap observed in water-exposed MAPI-PbAc2 films (Fig.5.2c) and the constant WF value for MAPI-PbI2 films (Fig.5.2d). The Pb:I elemental ratio for both materials (1:3) remains constant even after exposure to 100 cycles of water vapor, indicating that a possible degradation cannot be attributed here to perovskite segregation at the surface and thus a change in stoichiometry. Interestingly, the traps expected to be present in the MAPI-PbAc2 surface, VI and MAi, are typically positively charged, which would imply that they are keen to interact with the electronegative oxygen atoms from water molecules. As for MAPI-PbI2, the considered defects, VPb and Ii are negatively charged, which would make them more difficult for the oxygen to interact.

Figure 5 . 4 .

 54 Figure 5.4. (a) XRD diffractograms of the MAPI-PbAc2 pristine and water-exposed (b) MAPI-PbAc2 lattice distance displacement upon water vapor exposure, (c) XRD diffractogram of the MAPI-PbI2 pristine sample and (d) MAPI-PbI2 lattice distance displacement upon water vapor exposure. Lattice a (c) stands for the shorter (longer) axis in the tetragonal structure.

Figure 5 . 5 .

 55 Figure 5.5. XRD diffractograms of the MAPI-PbI2 pristine and water-exposed samples.

Figure 5 . 6 .

 56 Figure 5.6. AFM height images of a MAPI-PbAc2 film upon exposure to (a) 40 and (b) 150 cycles of water. Power spectrum density plot of (c) MAPI-PbAc2 and (d) MAPI-PbI2 films.

  7 and 143.3 eV for the Pb 4f7/2 or Pb3O4 and Pb 4f5/2 or PbO2 respectively, for both MAPbI3 perovskite films, Fig. 5.8c and f. However, there is a slight difference between both perovskite surfaces. While for MAPI-PbAc2 only the peaks attributed to Pb 4f can be found, Fig. 5.8c, in MAPI-PbI2

  7 and 631.0 eV for I 3d5/2 and I 3d7/2, respectively, Fig. 5.8a and d. Finally, the binding energy of the N 1s main peak (at 402.8 eV) remains unaltered as well, Fig. 5.8b and e.

Figure 5 . 7 .

 57 Figure 5.7. C1s core level XPS spectra of MAPI-PbAc2 (a,c,e) and MAPI-PbI2 (b,d,f): pristine (a, b), exposed to 40 cycles (c, d) and to 100 cycles of water (e, f).

Figure 5 . 8 .

 58 Figure 5.8. XPS core level spectra of (a) (d) I 3d, (b) (e) N 1s and (c) (f) Pb 4f for pristine and water-exposed MAPI-PbAc2 samples. (a-c) Correspond to MAPI-PbAc2 and (d-f) to MAPI-PbI2. Blue curves represent the pristine film, orange represents the film exposed to 40 cycles and green represents an exposure to 100 water cycles.

Table 1 :

 1 Phase transitions of MAPbI3 determined from synchrotron X-ray powder diffraction data. Temperatures were equilibrated for approximately 10 minutes after 1 K changes.25

Table 1 :

 1 Materials used for the preparation of MAPI-PbAc2 and MAPI-PbI2 perovskites. leading to a final thickness of 5 nm. The chamber was set to 200 ºC and the chuck at around 55 ºC. The pressure in the vacuum chamber was 1.1 torr.

	MAPI-PbAc2 route	MAPI-PbI2 route
	MAI (Greatcell Solar)	MAI (Greatcell Solar)
	PbAc2•H2O	PbI2 (99% Merck)
	(99.999% Merck)
	HPA (50% Merck)	DMF (Merck)
	DMF (Merck)	DMSO (Merck)
		Ethyl Acetate (Merck)

Additionally in Chapter 4, perovskite devices were characterized with an AFM exposed to air. Therefore, an insulating layer of aluminium oxide (Al2O3) was deposited on top of the perovskite layer. Using an atomic layer deposition (ALD) chamber, trimethylaluminium (Al(CH3)3, 97% Sigma Aldrich) and water were injected in a series of pulses,

  2.18a. As for the C-V measurements, we found that taking multiple images of the same 2 x 2 µm spot and, varying the applied voltage sent to the tip, was the best method to obtain reliable capacitance measurements. The S11 values from the images taken at each bias voltage were averaged. Then the S11 values were converted to admittance values following Eq. 2.34, and then converted to capacitance values by taking the imaginary part of the admittance. Although this calibration model does not provide quantitative results, it is still useful to qualitatively compare the local capacitance. First, the MAPI-PbAc2 perovskite layer was tested, which topography and derived capacitance can be seen in

  This discrepancy in activation energies and diffusion ions make it hard to draw any conclusion from our results. In addition, in our case, with a device structure such as PEDOT:PSS/MAPbI3/PtIr, the activation energies may be also different, thus, it remains an open question whether the high hysteresis in C-V curves is caused by Ii or by other species. As discussed in Chapter 1, hysteresis can be induced by several phenomena, and, from the I-V measurements from Fig.4.9b and with the before mentioned literature, we could think about mobile ions inducing a hysteretic capacitance, however, at this point we cannot discard other phenomena. In order to confirm mobile ions or defects as the cause of the hysteresis, a more detailed and exhaustive study of this material should be done. For instance, impedance spectroscopy measurements may help to detect and/or compare the capacitance with the capacitance measured with the SMM. Also, DFT calculations on the activation energies for defects in our device architecture would point towards specific defects, therefore, a specific study could be done. Nonetheless, it has to be pointed that SMM measurements would be an interesting technique to detect hysteretic properties at the local scale when combining it with other studies.

  MAPI-PbAc2 and MAPI-PbI2 perovskite films, previously described in Chapter 3, were exposed to water vapor in this chapter. First, the perovskite films were introduced in a low-vacuum chamber (~ 1.15 torr) with 0.88 dm 3 of volume. Water vapor was then injected by cycles into the chamber, heated at ~ 55 ºC, and purged by a continuous nitrogen flow to hinder any condensation. Each cycle corresponds to a pulse of water and then flushing it away before the next one. This method allows a very precise control of the water quantity to which the sample is exposed. For comparison, Li et al. reported a threshold of 2•10 10 L (1 L corresponds to 10 -6 torr.s) for MAPbI3 to decompose, based on XPS measurements.8 In this work, samples were exposed to much lower amounts of water, i.e., from 2.25•10 8 L up to 5•10 9 L, which allows to follow earlier stages of the degradation process.

Acknowledgements

The results presented in this thesis would not be possible without the help and support from the group of brilliant scientists and people I had the fortune to share my PhD journey.

I will name some of the people that influenced me the most through these years.

Chapter 3

Local Electronic Heterogeneities on the CH 3 NH 3 PbI 3 Perovskite Surface

The following publication was adapted from this chapter: J. Llacer, D. Moerman, O. Douhéret, X Noirfalise, C. Quarti, R. Lazzaroni, D. Théron and P. Leclère, ACS Applied Nano Materials, 2020, 3, 8268-8277.

Surface properties are a key aspect to improve not only the interface properties in photovoltaic cells but also the stability. They are even more important in perovskite materials where much engineering has been researched in order to improve the interface properties. In this chapter, the local electronic properties of two different CH3NH3PbI3 perovskite films are studied by atomic force microscopy methods. Nanoscale features such as local surface potential are correlated to current response. Perovskites made with lead acetate as a precursor for lead, result in a more heterogeneous surface compared to perovskites made with lead iodide in the precursor solution.

Topography Crosstalk

For this section, 5 nm of Al2O3 were deposited on top of the perovskite layer (see section 2.1). The SMM was operated in ambient conditions so the insulating layer prevented the moisture interacting with the perovskite surface while measuring the electronic local properties. Since the operating SMM frequency influences the depth of analysis, all the SMM measurements were carried out at 18.1 GHz. At such high frequency, the wave penetration is weaker so the measurement remains located at the surface. 5 Surprisingly, the capacitance variation with bias voltage shows a more pronounced hysteresis compared to MAPI-PbAc2, Fig. 4.9a. In order to confirm the hysteresis effect, current vs voltage images were recorded as in Fig. 4.9b, but when looking to the current variation it does not seem to be as much hysteretic. Nonetheless, MAPI-PbI2 is expected to be n-type (as explained in Chapter 3), but here we found positive current for positive voltages. This was previously studied by Bowring et al. and they found that mobile ions accumulate at the contact and are responsible for the current flowing in reverse bias, this may result in a decrease in efficiency. 9 Defects can be also considered as mobile ions, from Chapter 3 we know MAPI-PbI2 has an n-type character, thus, the defects are most probably either VPb and/or Ii. For the p-type MAPI-PbAc2 perovskite surface, the most probable defects to be present are either VI and/or MAi. In fact, Ii are believed to be the main source of trapping charges in MAPbI3. 10,11 Yang et al. found that VI has lower formation energy than MAi and Ii and also are fast diffusion defects. Therefore, they concluded that VI should be the primary contributor to ionic conductivity in MAPbI3. 12 Nonetheless, VI can show different activation energies depending on the device architecture. For instance, while Eames et al. found 0.6 eV for the Ii activation energy in TiO2/MAPbI3/Spiro-OMeTAD/Au, Bryant et al. found that it is takes only 0.12 eV for Ii to be active in PEDOT:PSS/MAPbI3/PCBM/LiF/Ag. 13,14 On the other hand, Yang et al. found a high ion conduction caused by iodide ions, which could lead us to thing that it is rather Ii the defects causing the high hysteresis. Stability is one of the most studied properties in perovskite materials. In this Chapter, stability against moisture degradation is studied at the nanoscale. The different regions found in Chapter 3 are also analyzed in this Chapter upon small amounts of water exposure. By controlling the amount of water at which the samples are exposed, a higher stability was found for perovskites prepared with lead iodide in the precursor solution.

The results confirm the presence of different surface defects, as previously suggested in Chapter 3. and 28.31º, corresponding to the X-ray scattering from the (110) and (220) planes, Fig. 5.4a. 18 After exposing the MAPI-PbAc2 sample to 100 cycles of water, there is a reduction in the X-ray scattering intensities while the characteristic PbI2 signal is not present, indicating that the perovskite did not start structural degradation yet. On the other hand, in the MAPI-PbI2 film, apart from the most intense peaks at 14.07º and 28.40º, there is another intense peak at 31.82º attributed to the (222) plane, Fig. 5.4c. The low-intensity feature at 12.65º in the XRD diffractogram for MAPI-PbI2 indicates the presence of unreacted PbI2. 19 The reduction in intensity of the XRD features seen for MAPI-PbAc2 could be related to the change in morphology observed with AFM, (Fig. 5.6b). Fig. 5.6c

shows the variations of the topographical roughness of the MAPI-PbAc2 layer upon exposure to water vapor, as described by the power density spectrum of the AFM images.

The gradual increase of the spectral density at high wavenumbers (i.e., smaller spatial distances) for MAPI-PbAc2 indicates that the relative roughness increases when increasing the water exposure, probably due to an increase of small grains. This is illustrated by the image of Fig. 5.6b, which shows the surface morphology after 150 cycles of water exposure (which corresponds to amounts of water similar to what reported by literature). The large grains initially present have broken up into smaller features, which is expected to strongly impact the electronic properties, as amply reported in literature. 8 It is also important to notice that the spectral density of the MAPI-PbI2 material is much less sensitive to water exposure (Fig. 5.6d), consistent with the better preservation of the electronic properties.

The highest XRD intensity peak in both MAPbI3 films is the (110) peak and it is used to measure the a lattice constant, which corresponds to the short lattice distance in a tetragonal geometry. The long lattice distance, the c lattice, is calculated mainly from the (222) peak. As a result, Fig. 5.4b and 5.4d show the evolution of the a and c lattice distances upon water exposure in MAPI-PbAc2 and MAPI-PbAc2 respectively. For MAPI-PbAc2, only the a lattice parameter is increased when exposing the sample to water. This lattice expansion could be related to a reaction between the MA + cation and water molecules, which has been proposed by Li et al.. They propose that water molecules act as n-dopants and move the perovskite Fermi level closer to the conduction band. This is discussed below on the basis of the XPS results and is consistent with the positivelycharged defects reacting with water. 8 In contrast, in MAPI-PbI2 both lattice distances remain constant up to 200 cycles of water exposure (Fig. 5.4d).

Chapter 6

Conclusions and Outlook

along with the solar cell performance could be addressed. Ideally, this would be achieved with both, the techniques performed in this thesis and DFT calculations. Regarding the SMM measurements, there is still range to improve the calibration procedure presented in Chapter 4. Although the calibration method that does not require any additional calibration sample, a deeper analysis would be needed in order to find how to extract real capacitive values. Nonetheless, the heterogeneous contrast in capacitance is still valid and it would suggest that point defects are present, as it has been demonstrated by KPFM and C-AFM. A way to corroborate this assumption would be to correlate SMM with KPFM or C-AFM by either installing a camera with a higher resolution in the AFM system using the SMM or using the IEMN-made SMM, which is paired with an SEM.