Hybrid and Inorganic Perovskite Nanostructures 9783110602173, 9783110601220

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Lioz Etgar Hybrid and Inorganic Perovskite Nanostructures

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Lioz Etgar

Hybrid and Inorganic Perovskite Nanostructures

Author Prof. Lioz Etgar Institute of Chemistry The Hebrew University of Jerusalem 91904 Jerusalem Israel [email protected]

ISBN 978-3-11-060122-0 e-ISBN (PDF) 978-3-11-060217-3 e-ISBN (EPUB) 978-3-11-060033-9 Library of Congress Control Number: 2020931694 Bibliographic information published by the Deutsche Nationalbibliothek The Deutsche Nationalbibliothek lists this publication in the Deutsche Nationalbibliografie; detailed bibliographic data are available on the Internet at http://dnb.dnb.de. © 2020 Walter de Gruyter GmbH, Berlin/Boston Typesetting: Integra Software Services Pvt. Ltd. Printing and binding: CPI books GmbH, Leck www.degruyter.com

Contents Abstract

1

1

Introduction

2 2.1

All-inorganic perovskite nanostructures 4 Kinetics – how to control the size distribution of Cs-based perovskite NPs 4 4 Synthesis of CsPbBr3 NPs 4 Tracking the size distribution of CsPbBr3 NPs 7 Synthesis of CsPbI3 NPs 7 Tracking the size distribution of CsPbI3 NPs Discussion 8 Ostwald ripening 10 Summary 11 Tuning the length and optical properties of perovskite NWs 12 12 Synthesis of CsPbX3 NWs Addition of hydrohalic acid (HX, X = Cl, Br, I) 14 Discussion 18 Summary 19 Changing the A site in perovskite nanostructures 20 Rubidium lead chloride nanocrystals: synthesis and characterization 20 The role of ligands 24 Summary 27 Near-ultraviolet to mid-visible bandgap tuning of mixed-cation 28 RbxCs1–xPbX3 (X = Cl or Br) perovskite NPs Summary 36

2.1.1 2.1.2 2.1.3 2.1.4 2.1.5 2.1.6 2.1.7 2.2 2.2.1 2.2.2 2.2.3 2.2.4 2.3 2.3.1 2.3.2 2.3.3 2.4 2.4.1 3 3.1 3.1.1 3.1.2 3.2

2

Hybrid perovskite nanostructures 37 Two-dimensional hybrid perovskite nanorods Role of the ligands 40 Summary 42 The effect of the alkylammonium ligands length on OIP NPs 42

37

VI

Contents

3.2.1 3.2.2 4

Study of the VDW interactions Summary 49 Summary and outlook

Acknowledgments References

53

51

50

45

Abstract Organic–inorganic perovskite and all-inorganic perovskite are semiconductor materials with exciting properties. In the last several years, these materials have been proved to be efficient light harvesters in photovoltaic solar cells and other optoelectronic applications. However, even the more fascinating research area related to these hybrid perovskites is their synthesis and properties in the nanoscale regime. Perovskite nanostructures can be synthesized by the metastases reaction using several techniques, while their optical properties can be tuned by halide or cation exchange and not necessarily by quantum confinement. This book aims at providing the tip of the iceberg to a whole new family of nanoscale semiconductor materials with more research and applications. This book provides an overview on hybrid and all-inorganic perovskite nanostructures. The role of ligands in these nanostructures and their influence on the shape and optical properties are discussed. Moreover, the replacement of the “A” cation in these perovskite nanostructures is demonstrated.

https://doi.org/10.1515/9783110602173-001

1 Introduction Lead halide perovskite materials are being studied intensively in the past years for their outstanding photovoltaic (PV) activity. Recently, the nanosize of lead halide perovskite evolves as a new independent field in the nanoscale materials. The formula for perovskite is AMX3. Perovskite enables high diversity from organic–inorganic hybrids to all-inorganic perovskites, mostly known as methylammonium (MA) lead halide (MAPbX3; X = Cl, Br, I) or cesium lead halide (CsPbX3; X = Cl, Br, I) compositions. Due to the diverse nature of perovskites, some advanced perovskite materials were studied using chemical modifications. There are reports on varying the monovalent cation [A = CH3NH3, formamidinium (FA) CH3(NH2)2, or Cs+], the divalent metal cation (M = Pb2+, Sn2+, and Ge2+), and the halide (X = Cl−, Br−, or I−) [1–13]. The opportunity to modify the perovskite chemically is of great importance. It enables to tune the bandgap and to adjust the properties to a specific application. Moreover, it opens the possibility to avoid toxic elements and to improve the synthetic routes. Nowadays, in the PV activity, the perovskite-based solar cells achieve more than 25% efficiency using a mixed-cation perovskite composition, which emphasizes the importance of chemically modified perovskites [14, 15]. As mentioned earlier, in the nanoscale regime, the metal halide perovskite is still behind. There are several reports on mixed-halide systems in the nanoscale [16–18], while mixed-cation systems are still behind both hybrid organic–inorganic and all-inorganic perovskite nanoparticles (NPs). Protesescue et al. reported on FA mixed with Cs in perovskite NPs with near-infrared emission [19], while a mixed metal cation system CsPbxMn1–xCl3 was reported with some new perspectives for tuning the optical properties of perovskite NPs [9]. The Goldschmidt tolerance factor (TF) is being used to predict whether a perovskite structure will be formed. It is important to consider that the optical properties of NPs are affected mainly by the electronic structure, while structural distortions are also known to influence and can be predicted by the TF. The TF is aimed at predicting a stable perovskite structure related to a three-dimensional (3D) cubic close packed ions [20–22]. The TF is calculated as p follows: t = ðrA + rX Þ= 2ðrM + rX Þ, where r is the ionic radius, and the empirical formability range is 0.8 < t < 1.0 [23]. Looking on the shape of the nanostructures following the first report on a synthesis of CsPbX3 nanocubes [16] very quickly, the field of CsPbX3 perovskite NPs emerged extensively and few groups started developing new methods to synthesize these particles. https://doi.org/10.1515/9783110602173-002

1 Introduction

3

The synthesis of CsPbX3 nanowires (NWs) was carried out by modifying the original synthesis. Long NWs were formed during longer growth time before quenching the reaction. Other reports show CsPbX3 nanoplates, where the injection temperature was lower than the one used for the original synthesis of nanocubes [24, 25]. The shape of NPs can be affected by the temperature of the synthesis, for example, at lower temperatures quasi-2D (two-dimensional) asymmetric morphologies were observed. Akkerman et al. [11] reported on the relation between the reaction temperature and the product’s shape. An important outcome of these studies is that the resulting shape and size of NPs are strongly dependent on the injection temperature and growth phase. However, not only the reaction temperature influences the shape and size of NPs, but the ligands also play a major role. In the synthesis of hybrid organic–inorganic perovskite (OIP) NPs, usually the solvent/anti-solvent combination is used as the trigger to form the NPs. Vybornyi et al. and Protesescu et al. reported the formation of MAPbX3 and FA lead bromide (PbBr2) cubic nanocrystals (NCs) using a hot injection approach [26, 27]. This approach requires high temperature and an inert-gas atmosphere, which complicate its utilization. In addition, those MAPbX3 NPs are unstable due to labile surface passivation, which allows fusion processes. The ligands in hybrid OIP NPs are alkylammonium cations [3, 6, 28]. Due to their positively charged ammonium group, these cations stabilize the surface of the perovskite NPs, specifically on the (00n) facets [29]. These facets comprise the exposed, negatively charged halides that electrostatically attract the ammonium cations [30]. The long aliphatic cation of the alkylammonium ligand does not fit into the confined octahedral hole of the perovskite structure and therefore inhabits the NP growth and passivates its surface. These alkylammonium ligands affect the electronic structure of the NPs (as discussed in the book); in contrast, the “standard” ligands stabilize the NP structure but usually do not affect the electronic structure [31, 32].

2 All-inorganic perovskite nanostructures 2.1 Kinetics – how to control the size distribution of Cs-based perovskite NPs In this chapter, we describe the investigation of growth duration on the perovskite NPs size distribution. To accomplish this, we synthesized CsPbX3 NPs, where X = Br or I. Four different growth durations were studied 1, 4, 20, and 40 s. We study these growth durations based on the previous observation [16] that the majority of growth of the CsPbX3 NPs occurs within the first 1–3 s, such that we could analyze the distribution focusing and defocusing beyond the main growth stage, between 1 and 40 s.

2.1.1 Synthesis of CsPbBr3 NPs We synthesize CsPbBr3 NPs over four different growth durations: 1, 4, 20, and 40 s. All products exhibit similar absorbance onset in the range of 515–522 nm and sharp photoluminescence (PL) peak in the range of 515–525 nm (depends on their size) with full width half maximum (FWHM) of 20 nm and PL quantum yield (PL-QY) of 60%. The absorbance and PL of CsPbBr3 NPs with growth duration of 4 s compared to the absorbance of bulk CsPbBr3 deposited by a one-step process [33] on the substrate are shown in Figure 1(a). The absorbance of NPs and bulk is similar in shape with slightly redshifted spectra of the bulk CsPbBr3 due to weak quantum confinement of the CsPbBr3 NPs. The electron diffraction (ED) pattern of the CsPbBr3 NPs is shown in Figure 1(b) and the corresponding diffraction-derived d-spacing is shown in Figure 2(a). The d-spacing matches the cubic structure of CsPbBr3. The XRD (X-ray diffraction) of the CsPbBr3 NPs presented in Figure 2(b) further supports the cubic structure.

2.1.2 Tracking the size distribution of CsPbBr3 NPs To follow the size distribution of NPs, the synthesis parameters of CsPbBr3 NPs were held constant except the growth duration variations, which varied between 1 and 40 s as shown in Figure 3. The different growth durations yielded cubic CsPbBr3 NPs with an average side length of 9.0 nm for 1 s growth, 9.6 nm for 4 s https://doi.org/10.1515/9783110602173-003

2.1 Kinetics – how to control the size distribution of Cs-based perovskite NPs

(a)

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Abs CsPbBr3 NPs PL CsPbBr3 NPs Abs Bulk CsPbBr3

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Figure 1: (a) Absorption and photoluminescence spectra of CsPbBr3, bulk, and NPs at 4 s growth time. (b) Electron diffraction pattern of CsPbBr3 NPs [34].

(a)

(b)

Figure 2: (a) d-Spacing of CsPbBr3 NPs obtained from electron diffraction and corresponding crystal planes. (b) X-ray diffraction spectra of CsPbBr3 [34].

growth, 8.4 nm for 20 s growth, and 9.3 nm for 40 s growth. High-resolution transmission electron microscopic (HR-TEM) measurements are observed in Figure 3(a)–(d), presenting the CsPbBr3 NPs of different growth durations and the corresponding size distribution histograms (Figure 3(e)–(h)). (The size distribution histograms were calculated from data obtained using ImageJ software, and the typical analysis population being 2,000 NPs.) It is shown that the FWHM of the size distributions increase with the growth duration, implying a defocusing process (Ostwald ripening), as will be discussed below.

(f)

(e)

(g)

(c)

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Figure 3: HR-TEM images of CsPbBr3 NPs allowed to grow for (a) 1 s, (b) 4 s, (c) 20 s, and (d) 40 s. Inset: HR-TEM of a single NP. normalized size distribution histograms for NP populations of each growth time; (e) 1 s, (f) 4 s, (g) 20 s, and (h) 40 s. FWHM of the size distributions is indicated [34].

(b)

(a)

6 2 All-inorganic perovskite nanostructures

2.1 Kinetics – how to control the size distribution of Cs-based perovskite NPs

7

2.1.3 Synthesis of CsPbI3 NPs To further study the variation in size distribution of the CsPbX3 NPs, we synthesized CsPbI3 NPs over four different growth durations: 1, 4, 20, and 40 s. All products exhibit similar absorbance onset in the range of 695–705 nm and sharp PL peak in the range of 684–695 nm (based on their size) with FWHM of 37 nm and PL-QY in the range of 75–77%. Figures 4 and 5 show detailed characterization of the CsPbI3 NPs with growth duration of 4 s. The absorbance and PL are shown in Figure 4(a) and HR-TEM is shown in Figure 4(b). The lattice fringes of the CsPbI3 NP are observed in Figure 4(b), corresponding to the (1,1,0) crystallographic plane. The cubic structure of the CsPbI3 NPs is supported by the ED (inset Figure 4(b)) and its corresponding d-spacing (Figure 5(a)) and the XRD (Figure 5(b)). (b) 1.0

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10 nm 400

500 600 700 Wavelength (nm)

Figure 4: (a) Absorption and photoluminescence spectra of CsPbI3 NPs. (b) HR-TEM image of a particle with distinct lattice fringes and inset of electron diffraction pattern [34].

2.1.4 Tracking the size distribution of CsPbI3 NPs As indicated previously, here we follow the growth duration variations between 1 and 40 s while all other synthetic parameters were held constant (Figure 6). The different growth durations yielded cubic CsPbI3 NPs with an average side length of 13.40 nm for 1 s growth, 10.5 nm for 4 s growth, 7.5 nm for 20 s growth, and 13.46 nm for 40 s. The size distribution was analyzed by HR-TEM measurements as shown in Figure 6(a)–(d), presenting the CsPbI3 NPs at different growth durations and the corresponding size distribution histograms (Figure 6(e)–(h)). (Analysis was performed using ImageJ software.) It can be

8

2 All-inorganic perovskite nanostructures

(a)

(b) 70

(200)

Intensity

60 50 40 30

(110) (221) (111) (210) (220) (211) (100) (222)

20 10 10

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30 40 2θ (deg)

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Figure 5: (a) d-Spacing of CsPbI3 and corresponding crystal planes; (b) X-ray diffraction spectra of CsPbI3; the XRD peaks match the crystallographic planes reported in Ref. [35, 36]. Taken with permission from Ref. [34].

seen that the FWHM decreases over 4 and 20 s of growth, and increases again over 40 s of growth, implying defocusing followed by focusing, followed by focusing, followed by defocusing as discussed below.

2.1.5 Discussion The size distribution behavior of CsPbBr3 and CsPbI3 varies. In the case of CsPbBr3, we observe a widening of the size distribution between t = 1 s and t = 40 s of growth duration; however, in the case of CsPbI3, we observe first a narrowing of the size distribution between 1 and 20 s growth duration, and then a widening of the size distribution between 20 and 40 s. Initially, particles smaller than the critical radius dissolve into the existing pool of monomers, while particles above the critical radius grow with a rate depending on their size: particles only slightly larger than the critical radius grow faster, while ones relatively larger grow slower. Since particles close to the critical radius grow fastest, they “catch up” to the larger particles, bringing more of the particle population into the same size range, narrowing the distribution [37]. The size-dependent growth rate is depicted graphically in Figure 7. Therefore, it is expected that a burst of nucleation is followed first by a phase of distribution narrowing (focusing), and then by a second phase of distribution widening (defocusing). The nucleation burst results in most particles being of similar size in an environment of high monomer concentration. (Monomer concentration is important because it is inversely related to the critical radius.)

(f)

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Figure 6: HR-TEM images of CsPbI3 NPs with growth time of (a) 1 s, (b) 4 s, (c) 20 s, and (d) 40 s. Inset: HR-TEM of single CsPbI3 NP at different growth times. Normalized size distribution histograms for NP populations of each growth time (e) 1 s, (f) 4 s, (g) 20 s, and (h) 40 s. FWHM of the size distributions is indicated [34].

(b)

(a)

2.1 Kinetics – how to control the size distribution of Cs-based perovskite NPs

9

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2 All-inorganic perovskite nanostructures

Figure 7: The size-dependent growth rate. r* – critical radius. Reprinted with permission from reference.

2.1.6 Ostwald ripening When monomer concentration is depleted as a result of growth, the critical radius increases past the average size. In that case, a large fraction of the population is below the critical radius and when dissolves, their monomers add onto larger growing particles. During focusing, most of the particles grow at a rate that allowed them to catch up in size to the larger particles. During defocusing, most of the particles shrink/disappear while some grow, broadening the distribution. The formation of small particles is kinetically favored, while large particles are thermodynamically favored. However, small particles have larger surface-to-volume ratio than large particles. When looking at the molecules on the surface of the particles, they are energetically less stable than the molecules in the interior of the particles. Therefore, small particles will attain a lower energy state if they will be transformed to large particles as in the case of Ostwald ripening. Figure 8 shows the focusing and defocusing processes. In the case of CsPbI3, we observed a moment of focusing at t = 20 s growth duration (stage 2 in Figure 6), and subsequent defocusing at t = 40 s growth duration (stage 4 in Figure 8). In the case of CsPbBr3, only distribution widening (defocusing) is observed (stages 3 and 4 in Figure 8). There is a difference in the focusing/defocusing processes between these two perovskites (e.g., CsPbI3 and CsPbBr3). This difference is related to the monomer concentrations. As stated earlier, the monomer concentration is inversely proportional to the critical radius. We discussed earlier that the majority of growth of CsPbX3 NPs occurs within the first 1–3 s, such that the growth after the first 3 s is

2.1 Kinetics – how to control the size distribution of Cs-based perovskite NPs

11

Figure 8: Schematic illustration of the focusing and defocusing processes. At stage 1, the reaction flask has a large population of mode-sized particles relative to that of smaller and larger particles, and monomer concentration is high. Small particles below the critical radius dissolve into the pool of monomers while mode-sized particles grow fast, resulting in focusing (stage 2). When monomer concentration is depleted due to growth, the critical radius is increased and mode-sized particles begin to dissolve, while the small population of large particles continue to grow, resulting in defocusing (stage 3). Over time, particles below the critical radius progressively dissolve while large particles progressively grow, accentuating the positive skew (stage 4) [34].

mainly limited by the diffusion of the monomers. It was reported that the diffusion coefficient of Br is larger than the diffusion coefficient of I [38]. Moreover, the synthesis temperature of the CsPbBr3 NPs is higher (170 °C) than the synthesis temperature of CsPbI3 (145 °C), which also suggests faster diffusion of the monomers. (The reason for the different synthesis temperature is due to precipitation of PbI2 at higher temperatures.) Therefore, it can be assumed that after 4 s of growth, the monomer concentration is more depleted in the case of CsPbBr3 than in the case of CsPbI3, which increases the critical radius and results in defocusing of the population.

2.1.7 Summary In this chapter, we present some properties of CsPbX3 (where X = Br or I) NPs and discussed their kinetics by tracking the growth durations. The size distribution was analyzed by TEM, which allow to follow the focusing and defocusing of the NP size distribution. In the case of CsPbI3, focusing of the size distribution was observed after 20 s of growth duration, while further increasing the growth duration results in defocusing of the size distribution. In the case of CsPbBr3, no focusing of the size distribution was observed. The main reason for the difference in the behavior of the size distribution as a function of the growth duration is related to the monomer concentration, which is depleted faster in the case of CsPbBr3 than in the case of CsPbI3,

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increasing the critical radius and resulting in defocusing of the population. To conclude, narrow size distribution could be achieved by choosing the appropriate growth time in the case of CsPbI3, while higher monomer concentration is probably beneficial in the case of CsPbBr3 in order to have narrow size distribution. Understanding the kinetics of these attractive NPs is critical for future use of them in optoelectronic applications.

2.2 Tuning the length and optical properties of perovskite NWs In this chapter, we present the synthesis of CsPbX3 NWs, where X varies between Cl, Br, and I. We demonstrate few unit cell thickness of NWs and the ability to tune their optical properties and at the same time their length. Hydrohalic acids (HXs) were added during the synthesis and affect NWs in two ways: (i) HXs shorten the NWs and (ii) the halide in the HXs participate in halide exchange reactions [16], influencing their optical features. Detailed characterization includes the structural, optical, and physical properties of the NWs. In addition, we discuss the mechanism of formation of these NWs.

2.2.1 Synthesis of CsPbX3 NWs In order to synthesize CsPbX3 NWs, hot cesium-oleate (Cs-oleate) and PbBr2 precursor solutions were prepared for the source of Cs+, Pb2+, and Br− ions. Csoleate was synthesized under inert conditions using cesium carbonate (Cs2CO3), oleic acid (OAc), and 1-octadecene (ODE). PbBr2 precursor solution was prepared in dimethylformamide (DMF). The reaction was performed at room temperature by adding hot Cs-oleate to a vial that contained ODE, OAc, oleylamine (OLA), and a variable amount of HX acid (0–10 µL). PbBr2 precursor solution was added to the vial and after 10 s, acetone, which acted as an antisolvent, was swiftly added to the reaction mixture in order to quench the reaction, forming freestanding CsPbBr3 NWs. The basic CsPbBr3 NWs were synthesized as a first step without any addition of acid. Figure 9(a) and (b) shows the TEM images of the product from this synthesis. Micron-sized asymmetric nanostructures with width of ~3 nm are observed in this case. The obtained NWs were extremely narrow, while a width of ~3 nm corresponds to 5 unit cells [25]. In addition, the NWs show a narrow size distribution with an average width of ~3.3 nm.

2.2 Tuning the length and optical properties of perovskite NWs

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Figure 9: (a) and (b) Typical high-resolution transmission electron microscopy (HR-TEM) images of CsPbBr3 NWs without addition of any hydrohalic acid. The inset of (b) shows a photograph of the sample under UV light (λ = 365 nm). (c) Normalized absorption and PL spectra of CsPbBr3 NWs without addition of any hydrohalic acid. The solid line refers to the absorption spectrum and dashed line refers to the PL spectrum. (d) X-ray diffraction (XRD) pattern of CsPbBr3 NWs without addition of any hydrohalic acid corresponds to a standard orthorhombic pattern of CsPbBr3 [42].

The PL and absorbance spectra are shown in Figure 9(c), while the emission from CsPbBr3 NWs dispersed in hexane under UV light is observed at the inset of Figure 9(b). The PL of CsPbBr3 NWs is located around 475 nm and is blueshifted relative to the reported cubic CsPbBr3 NPs emission peak of ~519 nm, without the addition of HX in the precursor solution. The significant blueshift in this case is attributed to quantum confinement of the NPs. According to the effective mass theory [39], the calculated Bohr diameter of CsPbBr3 is 7 nm [7], which can explain the large blueshift. In addition, since the synthesis took place at room temperature, a better control on the growth can be achieved. XRD pattern of the synthesized CsPbBr3 NWs is presented in Figure 9(d). The CsPbBr3 NWs show the orthorhombic crystal structure of CsPbX3 [40], which is supported by the low-temperature synthesis,

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2 All-inorganic perovskite nanostructures

enhancing crystallization in the orthorhombic phase, as previously reported [41]. The (004) plane has high intensity and sharp peak, which indicate the asymmetric morphology of NWs.

2.2.2 Addition of hydrohalic acid (HX, X = Cl, Br, I) In the next step, for the synthesis of NWs, HX was added in a series by increasing the amounts as follows: 2.5, 5, 7.5, and 10 µL. Variable amounts of HX addition resulted in CsPbX3 NWs with different lengths. Figure 10(a)–(c) shows TEM images of the synthesized NWs with different HXs: HCl, HBr, and HI, respectively. (a)

(b)

(c)

Figure 10: TEM images of the synthesized CsPbX3 NWs. (a)–(c) A series of samples with increasing amounts of HCl, HBr, and HI, respectively. The arrows indicate the direction of increasing the amount of HX acid [42].

It can be observed that NWs retained their shape and morphology while increasing the HX amount. In each series of NWs (Figure 10(a)–(c)), the arrows indicate an increase in the amount of HX, while the length of the NWs was shortened with almost every increase of HX.

2.2 Tuning the length and optical properties of perovskite NWs

15

On the other hand, the width of NWs remained the same, ~3 nm. A slight difference in the tendency of shortening the length of NWs among the three HX-based NWs can be observed, which most probably arises from the acidity differences. When the acidity environment is high, it leads to an enhanced protonation of OLA ligands into oleyl-ammonium cations, randomly occupying binding sites of Cs+ ions in the lattice, and causing a shortening at a given amount of acid (discussed later in more details). This interruption to the crystal growth creates a passivation layer and limits further crystallization. As shown in Figure 10, the shortening effect is different among the acids. The strength of an acid depends on the equilibrium constant. HXs are considered strong acids, where HCl is defined as a strong acid and both HBr and HI are defined as very strong acids. This is based on the few factors, including the free energy of bond breaking and hydration [43, 44]. Based on this, it is expected that HI and HBr will have a stronger shortening effect in relation to HCl. However, this expectation does not necessarily fit the observations because of the use of different concentrations of the acids and examining the shortening effect of each acid separately. The study of various HXs is aimed at showing the tendency in which protons enhance shortening of the NWs and is consistent among different HXs. Figure 10(a) shows that the shape of the NWs changes upon adding more HCl, which can be explained by the acidity as discussed earlier. Moreover, it can be shown that more harsh acidic conditions may affect other facets of the NWs making its shape more plate-like. The random nature of the passivating effect by oleyl-ammonium cations is prominent in Figure 10. Each fraction of HX results in a large distribution of the length of NWs, which was hard to estimate. In the case of addition of HBr, there are less side products than in the case of the addition of HCl as shown in Figure 10(b) and (a), respectively. This may be related to the spontaneous halide exchange that had two identical Br− ions. This ion exchange is less destructive to the crystal structure than the ones with HCl and HI. Apparently, in Figure 10(c), the influence of HI is less dramatic, while the difference between every two consequent fractions (e.g., the fractions of 7.5 and 10 µL) is milder. The black dots that appear in some of the TEM images were already reported in the previous work [24] and were recognized as lead particles that form after an exposure to the electron beam. XRD of the synthesized NWs are observed in Figure 11(a). The effect of the different HX can be seen in the measurement. Addition of HCl and HI led to partial ion exchange processes, depending on the amount of acid added, while HBr showed no significant change in the crystal structure because it is an exchange of the same halide.

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(a)

(b)

Figure 11: (a) X-ray diffraction (XRD) patterns of CsPbBr3 NWs with different amounts of hydrohalic acids, including the orthorhombic standard pattern of CsPbBr3. Asterisk indicates nonrelated peaks of Cs4PbBr6. (b) Atomic force microscopy (AFM) scan of NWs deposited on freshly cleaved highly oriented pyrolytic graphite (HOPG). A NW of uniform 3 nm height is shown in the image. In accordance with TEM images, it is assumed to be a single NW. No lower NWs were imaged. Several NWs of similar height were observed in different scan areas as well as higher structures, assumed to be bundles. The inset represents a graph that corresponds to the scanned area [42].

2.2 Tuning the length and optical properties of perovskite NWs

17

For HCl-based NWs, there is a small right shift to higher angles, and the opposite trend is observed for HI-based NWs in relation to the peaks of orthorhombic CsPbBr3 crystal structure. The shift to higher angles can be explained by the halide size and the space it occupies in the lattice. Cl− is smaller than Br−; thus, the substitution of some of the Br− anions with Cl− anions can lead to a contraction of the lattice. Similarly, substitution of Br− anions with I− anions can lead to an expansion of the lattice [45]. Moreover, looking on the XRD spectra in Figure 11(a) an additional phase is observed, which belongs to the Cs4PbBr6 along with the CsPbBr3 phase. The Cs4PbBr6 phase was already reported [46]. The asterisks in Figure 11(a) indicate peaks that are attributed to the Cs4PbBr6 phase and not to the orthorhombic crystal structure of CsPbX3 perovskite. This phase was reduced in the case of HCl and increased in the case of HI, which might be due to mismatch or match of the crystal structure, respectively, between the exchanged crystal and the Cs4PbBr6 phase. In order to determine the height of NWs, we performed atomic force microscopic (AFM) measurement. The AFM scanning showed elongated structures of micrometers length, the lowest of which had a rather uniform height of ~3 nm (see the inset in Figure 11(b)). Probably, the 3 nm high structures are single NWs, in accordance with dimensions seen in TEM imaging. Higher structures may be bundles of several NWs or single NWs wrapped in organic substance. Optical characterizations are observed in Figure 12(a). The significant blueshift in this case is related to quantum confinement of the NWs, as mentioned earlier. In some cases, two peaks are recognized in the PL spectra, which indicate two possible population: one is NWs and the other might be cubic-shaped CsPbBr3 NPs that are redshifted in their emission. Due to the ionic nature of the perovskite crystal [16], a fast ion exchange can easily occur among halides in solution. Figure 12(a) shows an addition of either HCl or HI, which causes a blueshift or a redshift, respectively, in the absorption and emission spectra. As the ion exchange process proceeds, more ions are exchanged within the crystal, and the exciton peak shifts to higher or lower wavelengths, according to the exchanged lattice. In the case of HBr-based NWs, the optical properties are not changing due to the same halide involved in the exchange. The emission from NWs under UV light is shown in Figure 12(b)–(d). The different emission wavelengths are observed in these images due to the halide exchange. For HI-based NWs, the emission redshift from blue to green (Figure 12(b)) indicates that some bromide ions were exchanged with iodide ions, and the emission blueshift from light blue to dark blue (Figure 12(d)) indicates that some bromide ions were exchanged with chloride ions.

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(a)

(b)

(c)

(d)

Figure 12: (a) Normalized absorption and normalized photoluminescence (PL) spectra of the various samples of CsPbX3 NWs. The solid lines refer to absorption spectra and dashed lines refer to PL spectra. (b)–(d) Photographs of the synthesized NWs with addition of HI, HBr, and HCl acids, respectively, under UV light (λ = 365 nm) [42].

2.2.3 Discussion A schematic illustration describing the effect of HX on the length of NWs is shown in Figure 13. On the left of the figure, NWs that were synthesized without adding acid to the reaction mixture can be observed. Addition of HX increases the acidity of the reaction solution, resulting in the protonation of the OLA ligands from OLA into oleyl-ammonium cations that behave similarly to Cs+, as previously suggested by other groups as well [47, 48]. It is suggested that the OLA molecules are competing the Cs+ during the crystallization of CsPbX3. The competition is on the active surface of the growing NWs. The binding of oleyl-ammonium ligands creates a passivation layer on the growing surface of the NWs.

2.2 Tuning the length and optical properties of perovskite NWs

19

Figure 13: Schematic illustration of the passivation effect by hydrohalic acids on the length of CsPbX3 NWs [42].

In a typical synthesis, Cs+, Pb2+, X−, and oleyl-ammonium cation immediately crystallize into NW structures after an addition of a polar solvent (acetone was used here), which induces precipitation of NWs. In the current synthesis of NWs, some of the positions of the Cs+ at the growing surface are taken by the oleyl-ammonium cations, thus blocking further growth in the favored direction by the long oleic chain. The oleyl-ammonium ligands create a steric interference and probably cannot function as intrinsic organic cations within the perovskite crystal. The competition between Cs+ and oleyl-ammonium cations in addition to the amount of HX creates random short NWs and length passivation effect.

2.2.4 Summary In this section, we described the synthesis of Cs-based NWs and the possibility to control their length and optical properties by the addition of HX (X = Cl, Br, and I). The synthesized NWs show narrow size distribution of ~3 nm having two quantum-confined dimensions proved by TEM and AFM measurements. HXs influenced the NWs into two aspects: (i) shortening of the NWs by increasing the amount of acid added and (ii) alteration of compositional structure and optical properties induced by halide exchange reactions in the case of nonidentical ions of the acid and the crystal. For example, such a reaction can be induced by addition of HCl to the precursor solution that contained Br− ions. The acidity of HXs is a decisive factor for the properties of NWs, as well as the halide content in the acid. The HXs are considered as strong acids that can promote protonation of amines. The protonation of the amine ligands was the key to the shortening of micrometer-long NWs.

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The option to control the length of NWs and optical properties at the same time opens the possibility to use them in optoelectronic applications. Moreover, the asymmetric nature of the reported NWs paves the way to an oriented growth of another semiconductor via an epitaxial growth of two or more semiconducting materials. This produces systems for energy transport taking into consideration the optical and physical features of such nanocomposites.

2.3 Changing the A site in perovskite nanostructures The most common inorganic cation that is suitable to be at the “A” site in the perovskite structure is Cs+. The reason is due to its size, this is large enough to fit into the octahedral cage. In this section, we present a new possible cation that can be used in all-inorganic perovskites. The Rb+ cation is little bit smaller than the Cs+; therefore, according to the TF it cannot form a perovskite structure in the bulk. In this section, we show for the first time the synthesis and characterization of Rb-based NPs.

2.3.1 Rubidium lead chloride nanocrystals: synthesis and characterization The synthesis of rubidium lead halide NCs involved two precursor solutions. The first solution contains Rb2CO3, OAc, and ODE, and the second solution contains lead chloride (PbCl2), OAc, OLA, tri-n-octylphosphine (TOP), and ODE, while OAc and OLA were in a volume ratio of 1:1. The two flasks were heated to 120 °C for 1 h under vacuum conditions for degassing. Following degassing, the temperature was raised to 150 °C in both flasks under argon flow. The Rb-oleate precursor was injected into the PbCl2 precursor solution and the reaction was quenched with an ice bath after a few seconds. The product was precipitated twice with isopropanol to give a white precipitate and redispersed with hexane for further characterizations. The synthesis resulted in NCs with a square-like shape. Figure 14(a) shows HR-TEM images of the NCs. The inset of Figure 14(a) presents a magnified image of the NCs, revealing some clear black spots upon the NCs (discussed below in further details). The stability of these NCs is extremely high. Figure 14(b) presents the absorbance spectra of NCs measured at the day of the synthesis and after ~4 months while stored in ambient conditions. The absorbance peaks remain unchanged, showing the long-term stability of these NCs. The bandgap of these NCs is 4.05 eV, which is estimated from the absorbance spectra.

2.3 Changing the A site in perovskite nanostructures

(a)

(c)

21

(b)

(d)

(e)

Figure 14: (a) A high-resolution transmission electron microscopy (HR-TEM) image of rubidium lead chloride NCs. Inset: a magnified image of a few particles. (b) Absorbance spectra of the NCs at the day of the synthesis and after 4 months. Inset: A size distribution histogram of the NCs measured on 647 NCs. (c) The conduction band minimum (CBM) and the electron effective masses along G-M and G-Z crystal directions. (d) The same as (c) for the valence band top (VBT). (e) The dielectric function: its real part (left panel) and imaginary part (right panel), calculated along the crystal directions (red and black lines) [50].

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Therefore, the optical activity of the obtained NCs is in the UV spectral region. The bandgap of the bulk crystal obtained from (density functional theory) DFT calculations is 3.01 eV and from pseudopotential self-interaction correction method is 4.13 eV [49]. This is in good agreement with the absorbance measurement (not shown). The size distribution histogram based on more than 600 NCs is shown in the inset of Figure 14(b), evaluating the average side length of the NCs to be 11.29 ± 0.04 nm (the size distribution histogram was calculated from data obtained using ImageJ software). In this particle analysis, there was an assumption of square-shaped NCs. Figure 14(c) and (d) shows the band structure-derived effective masses, averaged over the crystal directions at the valence band top (VBT) and conduction band minimum (CBM), which are 0.53 and 0.19 of the electron mass (me) for holes and electrons, respectively. These numbers, together with the calculated dielectric function and its static value ε(0)/εvacc = 4.1 (Figure 14(e)), led to the exciton Bohr radius a0 = 1.584 nm and the exciton binding energy Eb = 0.111 eV. There is a close resemblance between the obtained parameters for Rb6Pb5Cl16 and the corresponding numbers for CsPbX3 compounds [2]. XRD measurement was performed and the results are presented in Figure 15(a). We found that the Rb6Pb5Cl16 tetragonal phase was observed by the XRD pattern, which is a non-perovskite phase. This phase was discovered by Monzel et al. [51] during their work on redetermination of the phase diagram of RbCl/PbCl2. One of the interesting insight of this work was that the perovskite phase RbPbCl3 cannot exist at room temperature, but as a mixture of the two previously unknown phases, Rb6Pb5Cl16 and RbPb2Cl5. Based on the phase diagram only with the application of heat, a phase transition can occur to get the perovskite phase. A temperature above 320 °C provides a tetragonal phase of RbPbCl3. Higher temperatures resulted in the cubic phase of RbPbCl3 perovskite. A repeating XRD measurement of the NCs detected 9.8% of Pb2O3 phase inside the Rb6Pb5Cl16 NCs and 90.2% of Rb6Pb5Cl16 phase. We can assign the Pb2O3 phase to the black spots in the Rb6Pb5Cl16 NCs shown in Figure 14(a). The are some reports on CsPbX3 NCs with black spots that appear as a response to the electron beam during the TEM analysis [25]. However, in this case the black spots were integrated in the NCs and were observed in “fresh” areas of the TEM grid, with minimal exposure to the beam. Nonetheless, the black spots became clearer and larger with longer exposure to the electron beam along with decomposition of the brighter parts. The difference in the TEM images between gray areas and black areas indicate elements with high electron density, which is the Pb concentration in the NCs (Figure 14(a)).

2.3 Changing the A site in perovskite nanostructures

23

(a)

(b)

(c)

(d)

Figure 15: (a) Experimental (black) and simulated (red) X-ray diffraction (XRD) patterns of rubidium lead chloride NCs that show the correspondence to a tetragonal Rb6Pb5Cl16 phase. (b) A scanning transmission electron microscopy (STEM) image. Insets: High magnification of a Pb-rich particle inside Rb6Pb5Cl16 NC with d-spacing that corresponds to plane (111) of metallic Pb and its corresponding FFT. (c) X-ray photoelectron spectroscopy (XPS) spectrum of the NCs measuring electronic states in the valence band. Inset: A magnified area of the graph near the valence band edge. (d) A schematic overview of the band diagram of the NCs [50].

Dark scanning transmission electron microscopy (STEM) and energydispersive X-ray spectroscopy (EDS) were applied to confirm the composition of NCs and the black spots inside the NCs, as shown in Figure 15(b). EDS measurement detected the atoms Rb, Pb, Cl, and O in the region of interest. Following quantification analysis, the atomic percentage of the atoms has been determined to be 17.40, 28.89, 37.62, and 16.07, respectively. We can assume that the ratio between Rb and Pb is ~1:1, according to the XRD; hence, it can be suggested that the remained Pb atoms are possibly bound to the detected O atoms to create a lead oxide phase or alternatively, metallic lead NPs, as previously reported for similar observations [47].

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The fast Fourier transform (FFT) of a black spot within the NC as shown at the inset of Figure 15(b) is related to plain (111) of metallic Pb. EDS analysis (not shown) indicated a higher content of Pb, reinforcing the assumption of the black Pb-rich particles within the Rb6Pb5Cl16 NCs to be either lead oxide or metallic Pb NPs. Since the surface of NCs is Pb-terminated, it opens the octahedral hole with Rb+ in the center, which is a sensitive absorption site for the oxygen atom from OAc, leading to termination of its molecules with the H-C=O group, instead of COOH. This prediction is consistent with the crystallographic data, which report 9.8% of Pb2O3 phase. The rubidium atom takes a role of a cation in the crystal; thus, its properties to bind oxygen are opposite to the electronrich elements of group VI (i.e., Se); for example, in PbSe, the PbxOy-type structures are not found [52]. Valence band (VB) XPS measurement and surface photovoltage (SPV) spectroscopy were performed in order to study the electronic properties of these Rb6Pb5Cl16 NCs more deeply (Figure 15(c) and (d)). The VB XPS measurement gives information about the density and occupancy of electronic states in the VB of the material, while the edge of the graph gives an evaluation of the VBT. The inset of Figure 15(c) presents the extracted VB value, 0.46 eV below the Fermi level. The work function of the NCs was measured using the SPV method and found to be 4.63 eV. An overview of the band diagram measured by absorbance, SPV and XPS, of the NCs is presented in Figure 15(d).

2.3.2 The role of ligands The effect of ligands on the NC formation and shape is studied by changing the ratio of OA:OLA in a volume ratio of 1:2 and 1:3 while keeping the total molar content of both ligands constant. The use of OAc and OLA pair is well known in the syntheses of perovskite NCs [32]. When both OAc and OLA are participating in the synthesis, there is a dynamic equilibrium of acid/base [53]. The OAc donates a proton to the OLA and turns into an oleate anion, as the OLA ligand is protonated and turns into an oleyl-ammonium cation [5]. HR-TEM images of 1:2 OAc:OLA NCs are shown in Figure 16(a), where few morphologies are observed. Large rectangles and long thin NWs consist of small black dots and distorted square-like phase, similar to NCs with 1:1 ratio (Figure 14(a)). Figure 16(b) presents the ratio of 1:3 OA:OLA. There are similar observations as the 1:2 OA:OLA NCs. In this case, the large rectangles do not appear. When synthesizing the NCs with OLA alone (without OA), no product was observed. Therefore, the formation of these NCs requires a minimal volume of OA

2.3 Changing the A site in perovskite nanostructures

(a)

(c)

25

(b)

(d)

Figure 16: HR-TEM images of obtained NCs with volume ratio of 1:2 (a) and 1:3 (b) of oleic acid (OAc) and oleylamine (OLA), respectively (OAc:OLA). (c) Absorbance spectra of NCs in both cases. (d) An XRD pattern of the NCs with 1:2 ratio of OA:OLA that corresponds to a main orthorhombic phase of Pb(OH)Cl. The spectrum of NCs with the ligand ratio of 1:3 OA:OLA is similar to the one presented in (d) [50].

in order to completely dissolve PbCl2. Therefore, OA has an important role in solubilizing the PbCl2 [2]. In the case of excess of OLA, the shape of NCs is affected based on the HRTEM. It was previously claimed that the acid–base pair has a definite impact on the morphology and optical properties of a similar system of perovskite NCs [53]. The new morphologies can be an indication that the binding of OLA to the surface is chemical and not physical, so variable amounts of OLA ligands can change the growth directions of NCs by blocking certain crystallographic facets. When the OLA concertation is much more than the OA concentration, then the protonation process is reduced, resulting in less protonated OLA species, that is, less oleyl-ammonium cations. These conditions may increase the coordinative

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bonding of the unprotonated OLA species with the Pb2+ cations on the surface, rather than the electrostatic interaction of the oleyl-ammonium cations with the Cl− anions in the lattice. Figure 16(a) and (b) shows elongated structures, which explain the above discussion. The absorbance measurement is shown in Figure 16(c), which presents an optical redshift, relative to the OA:OLA 1:1 case, with peak position at ~326 nm and onset at ~336 nm. The shift in the absorbance is unexpected, which can be due to a possible change in the chemical composition. XRD measurement was applied for the 1:2 product. Interestingly, the XRD spectrum detected a main orthorhombic Pb(OH)Cl phase (Figure 16(d)). The mineral Pb(OH)Cl was reported earlier [55–58]. In addition, Chen et al. reported on Pb(OH)Cl nanotubes that absorb in the UV region at 258 and 292 nm [55], which is similar to the absorbance of the obtained nanostructures. The high amount of unprotonated OLA ligands may interact with large area on the surface of NCs and avoid Rb-oleate species from participating in the crystallization of NCs, yielding another Rb-free crystal phase of Pb(OH)Cl. The XRD of the 1:2 product also recognized minor content of Pb2O3 and Rb6Pb5Cl16 phases. In comparison, the XRD of the 1:3 product detected only the Pb(OH)Cl phase. It can be summarized that in the case of an increase in the OLA content, the formation of Pb(OH)Cl enhances, repelling Rb-oleate species in the growth processes. Theoretical simulations were done in order to get more insight about this suggested mechanism. The interaction of the Pb-terminated (001) surface with the OAc molecules was simulated. A slightly shorter molecule than the OAc molecule was used for the simulation. Since the COOH group is the functional group, it is attached to the surface of NCs. Three plausible geometric configurations were assumed (Figure 17), and their atomic structures were optimized. Two of these reactions end up with the oxygen atom on top of Rb from a cavity of the triangular shape set by three Pb surface atoms. The formation energy is −2.88 eV when the acid molecule swaps the -OH termination into =O and -H, attached directly to the C atom which lost =O. The -OH group can also build the surface; if it takes the position between two Pb atoms, then the formation energy is −2.59 eV (see Figure 17). Before the reaction, no charge transfer was observed to the surface; thus, the band structure was not affected by proximity with the acid (see Figure 17). In contrast, the adsorption of O or OH affects the bandgap, as expected. Both cases cause the redshift; it is about 0.4 eV for the oxygen adsorption case, and it seems to be a multiexcitation between many impurity states for the surface reaction with -OH. However, the oscillator strengths need to be further investigated with more advanced methods to find which optical transitions are active.

2.3 Changing the A site in perovskite nanostructures

27

Figure 17: Adsorption sites for O and OH from the oleic acid and the corresponding formation energies [50].

The opposite case was also studied when syntheses were performed with ligand ratios of 2:1 OA:OLA and 3:1 OA:OLA (Figure 18(a) and (b)). A distorted shape was observed when compared with the 1:1 OA:OLA NCs (Figure 18(a)). The role of OAc ligands is to control the size to be more isotropic, as shown in Figure 18(a), which is in contrast to the case of OLA excess. In Figure 18(b), there is no observation for NCs (1:3 OAc:OLA ratio). The reason for that can be insufficient amounts of OLA, which can severely affect the stabilization of the formed NCs. The OLA concentration is too small; therefore, NCs cannot be formed due to less coordinative bonding of oleyl-ammonium ligands with the metal cations on the surface.

2.3.3 Summary This section presents the synthesis of rubidium PbCl2 NCs as a novel UV-active material with intriguing features and complicated structure.

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(a)

2 All-inorganic perovskite nanostructures

(b)

Figure 18: HR-TEM images of rubidium lead chloride NCs with volume ratios of (a) 2:1 and (b) 3:1 of oleic acid (OAc) and oleylamine (OLA) ligands [50].

Square-shaped NCs were observed by the HR-TEM, while the XRD determined the symmetry of tetragonal Rb6Pb5Cl16 phase. Theoretical calculations of electronic and structural properties correlated with the experimental data. Studying the role of ligands in these NCs reveals the formation of a main phase of Pb(OH)Cl with an optical redshift related to the square-shaped NCs. A competition between Rb cations and OLA on the surface can explain the resulting structures, while a larger amount of OLA resulted in a pure Pb(OH)Cl, without any Rb atoms in the lattice. Shu et al. [52] suggest that Pb(OH)Cl material can function as an anode for lithium-ion battaries, which can be very attractive in the nanometric scale as well.

2.4 Near-ultraviolet to mid-visible bandgap tuning of mixed-cation RbxCs1–xPbX3 (X = Cl or Br) perovskite NPs Following the previous section, we were motivated to shift the absorbance of NCs to longer wavelength. A possible way to do that is to have a mixed-cation composition inside the NCs, which means a mixing of Rb+:Cs+ as the A-cation. To do that it is required to calculate the TF of several perovskite compositions. In the x-axis of Figure 18(e), increasing Rb+ contents (Rb+:Cs+ as the A-cation) are presented according to the formula RbxCs1–xPbX3 (x-values = 0, 0.2, 0.4, 0.6, 0.8, 1; X = Cl, Br, I). We used the report by Shannon [59] to consider the coordination number of the ions. For cubic perovskite, the coordination number of the metal and the halide is 6 and for the cation is 12. In this work, it is assumed that the

2.4 Near-ultraviolet to mid-visible bandgap tuning of mixed-cation

29

distortion is reducing the symmetry of the structure and the coordination number of the cation as well, from 12 to 8. Thus, the used effective ionic radii are Cs+ (1.74 Å), Rb+ (1.61 Å), Pb2+ (1.19 Å), Cl− (1.81 Å), and Br− (1.96 Å). Based on these calculations we demonstrate the possibility to introduce Rb+ cation with Cs+ cation in RbxCs1–xPbX3 NCs using hot injection method. In short, Rb2CO3/Cs2CO3 or a mix of them were combined with Rb+/Cs+-oleate precursors in a three-neck flask along with OAc and ODE. For the lead halide (PbX2; X = Cl, Br) precursor, PbCl2 or PbBr2 were loaded in an additional threeneck flask along with OAc, OLA, and ODE. TOP was added just in the case of PbCl2 for a complete dissolution. Both flasks were degassed under vacuum, and then the temperature was raised to 150 °C for the reaction. A desired volume of Rb+/Cs+-oleate was swiftly injected to the PbX2 flask. After ~5 s, an ice bath was applied to quench the reaction. The crude NCs were purified through centrifugation, with isopropanol as an antisolvent for further characterizations. Figure 19 presents the absorption and PL spectra of the different mixedcation perovskite NPs with chloride (RbxCs1–xPbCl3, Figure 19(a) and (b)) and bromide (RbxCs1–xPbBr3, Figure 19(c) and (d)). There is a blueshift in the absorption for both chloride and bromide, where the x values correspond to 0, 0.2, 0.4, 0.6, and 0.8. In the case of chloride, there is a shift of 0.13 eV and for bromide 0.07 eV. The absorption spectra in Figure 19(a) and (c) confirm that at higher Rb+ content, the absorption onset shifts to shorter wavelength, while the PL peaks are shifted in the same trend (Figure 19(b) and (d)). We also tried the mixed-cation Rb+/Cs+ NCs with iodide as the halide; however, the synthesis product seems to be unstable, also showing no optical shift. According to the literature, bandgap tuning commonly carried out using mixedhalide systems; however, mixed-cation systems are also affecting the bandgap. Most of the reports about mixed-halide systems and halide exchange reactions showed a significant optical tuning of perovskite NCs that can be applied to various utilizations [16–18]. Tuning the bandgap by the halides is due to changing the energy level of the VB as a result of different energies of their p orbitals. On the other hand, substitution of the A-cation influences the bandgap indirectly through structural distortions. The Pb–X–Pb angle is influenced by the A-cation size. As this angle is more distorted (whether it is larger or smaller than the ideal 180° angle) and the octahedral tilting is larger, the bandgap will be changed. The overlap of the antibonding orbitals of the Pb2+ metal cation and the halide anion is deformed compared to the ideal cubic perovskite structure. The energies of the VB and the conduction band are shifting upward, the antibonding interaction is weaker; therefore, the Pb–X bonds become more stable [16, 33]. This explains

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(a)

(b)

(c)

(d)

(e)

Figure 19: (a,b) Absorbance and normalized photoluminescence (PL) spectra of RbxCs1–xPbCl3 NPs (x = 0, 0.2, 0.4, 0.6, 0.8). (c,d) Absorbance and normalized PL spectra of RbxCs1–xPbBr3 NPs (x = 0, 0.2, 0.4, 0.6, 0.8). (e) The calculated tolerance factors of RbxCs1–xPbX3 (X = Cl, Br, I) perovskites as a function of the Rb content (x) range from x = 0 to x = 1. The gray dashed line represents the bottom limit of the empirical formability range of perovskite (0.8) [60].

2.4 Near-ultraviolet to mid-visible bandgap tuning of mixed-cation

31

why the energy level of the VB shifts downward, overall widening the bandgap and shifting the absorption to shorter wavelengths. The synthesis of x = 1 (Rb+ alone as the A-cation) was done with Cl and Br. The Cl-based NCs resulted in another phase of Rb6Pb5Cl16 as described in the previous section. Also when synthesized with only Br as the halide, NCs were not formed. Previous reports show that the RbPbX3 phase can be stabilized only at elevated temperatures (above 320 °C), which explain the results [51, 61–63]. The PL-QY of the NCs shows relatively high values (Figure 20(a)), similar to the ones reported for CsPbX3 (X = Cl, Br) NCs. It is hard to distinguish whether there is a pronounced trend while increasing the ratio of Rb+:Cs+. Moreover, the PL-QY strongly depends on the purification process of the NPs and the conditions of each synthesis. Moreover, these results suggest that addition of Rb+ maintains the good PL-QY of the CsPbBr3 NPs.

(a)

(b)

Figure 20: (a) Photoluminescence quantum yield (PL-QY) of the NPs with Cl (purple) and Br (green). (b) Photographs of dispersions of x = 0.2 and x = 0.8 (left and right vials, respectively) samples under ultraviolet light (λ = 365 nm) showing the fluorescence for the extreme molar ratios of Br-based NPs [60].

The bright PL of the obtained Br-based NPs under UV light is presented in Figure 20(b), showing green and turquoise-blue emissions for low and high Rb+ contents, respectively. The changes in the emission wavelength are small; therefore, only extreme ratios of Rb+:Cs+ are presented in this picture. The change in the Rb+ contents in the case of Cl-based NCs is in the near-UV spectral range, so the change in the emission colors of the NCs with different Rb+ contents is indistinguishable by human eyes. The powder XRD (pXRD) of the RbxCs1–xPbX3 (x = 0, 0.2, 0.4, 0.6, 0.8, 1; X = Cl, Br) NCs is presented in Figure 21. In both cases of Cl and Br, the diffractograms

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(a)

(b)

Figure 21: Powder X-ray diffraction (PXRD) patterns for RbxCs1–xPbCl3 (x = 0–1) NPs. (b) PXRD patterns for RbxCs1–xPbBr3 (x = 0–0.8) NPs. Theoretical peak positions of orthorhombic and cubic (for comparison) CsPbCl3 (a), CsPbBr3 (b) and tetragonal Rb6Pb5Cl16 phase (top of (a)) are shown by vertical bars; on the right are zoomed fragments of XRD patterns [60].

show a perovskite crystal structure, where there was a low Rb+ contents. In the case of x = 0, there are peaks that correspond to a orthorhombic CsPbCl3 perovskite structure [13]. Higher Rb+ contents of x = 0.2 and x = 0.4 resulted in the same CsPbCl3 perovskite peaks. A slight shift of the perovskite peaks can be observed, which is associated with small changes of the values of the unit cell parameters upon Rb–Cs substitution. When the Rb+ content increases, the perovskite peaks become weaker and other peaks appeared. In the case of x = 0.6 and x = 0.8, peaks of RbPb2Cl5 phase were detected while at x = 0.8 peaks of orthorhombic perovskite are absent. Finally, in the case of x = 1 there are different peaks, which

2.4 Near-ultraviolet to mid-visible bandgap tuning of mixed-cation

33

correspond to a tetragonal Rb6Pb5Cl16 phase rather than perovskite RbPbCl3, as mentioned earlier. The synthesis of Rb6Pb5Cl16 was described in the previous section, which indicates the change in the crystal structure when replacing the Cs+ cation with the smaller Rb+ cation in the lattice, supporting the optical measurements that indicated mixed Rb+/Cs+ perovskite NCs. Therefore, it can be concluded that adding further the Rb+ content at the expense of Cs+ will cause geometrical changes, which disrupt the perovskite stabilization based on the cation radii. In Figure 21(b), the orthorhombic CsPbBr3 perovskite structure is obtained for x = 0. Low Rb+ contents yielded the same perovskite crystal structure, excluding the x = 0.2 case, in which two intense peaks are observed. These intense peaks may relate to impurities of Cs4PbBr6 phase that is characterized by peaks in the angles 12.8 and 25.9 that can correspond to the observed peaks. Peaks of Rb4PbBr6 and RbPb2Br5 phases are observed for x = 0.6 and x = 0.8 cases. The Rb4PbBr6 phase was previously obtained using solid-state reactions, starting from binary precursors [64]. It is possible that there is a presence of a small impurity of additional phases (Rb6Pb5Cl16, RbPb2Cl5, Rb2PbCl4, Rb3PbCl5, RbPb2Br5) in the synthesized material. Looking at the phase diagram reported by Monzel et al. [51], a 1:1 mixture (i.e., RbPbCl3) will not form a single perovskite phase at room temperature but instead it will form a two-phase mixture. In the reported work, the discussed phase was rhombohedral K4CdCl6-type structure in contrast to an earlier paper, which suggested a tetragonal Tl4HgBr6-type structure [65]. In this work, the pXRD results in a tetragonal Rb4PbBr6 structure for x = 0.6 and x = 0.8 products. In order to decide which phase will be formed is mainly stability related. It was concluded that only ns2-type of A-cations (such as In+ and Tl+) can stabilize the tetragonal Tl4HgBr6-type structure due to polarization effects and high electronegativity, compared with alkali ions of comparable size, such as Rb+. Possibly, a much “looser” crystal structure is formed in the nanoscale, which contributes to the formability of less preferred structure type [66]. To be more specific, when increasing the amount of Rb+ (over Cs), there are more phase modifications in both Cl- and Br-based NCs. The reason is due to the octahedral tilting; when the octahedral tilting is large, a more stable crystal phase is formed until the perovskite crystal structure is lost completely (e.g., in Cl; where x = 1). While in the case of Br, already at x = 0.8 the perovskite phase was not formed. The different NCs with a square shape are shown in Figure 22 obtained by TEM. We measured the size distributions of each product with different Rb+ content, using ImageJ software (not shown). It can be seen clearly that when increasing the Rb+ content in the NCs, the average size of the NCs is decreasing

34

2 All-inorganic perovskite nanostructures

(a)

(b)

Figure 22: (a) Transmission electron microscopic (TEM) images of RbxCs1–xPbCl3 (x = 0, 0.2, 0.4, 0.6, 0.8) NPs with the corresponding tolerance factors (TF). (b) TEM images of RbxCs1–xPbBr3 (x = 0, 0.2, 0.4, 0.6, 0.8) NPs with the corresponding TF. The scale bars correspond to 50 nm [60].

for both Cl and Br. When adding more rubidium to the crystal, at the expense of cesium, the d-spacing among the crystallographic plains is decreasing because Rb+ is smaller than Cs+, affecting the average size of the NCs with increasing Rb+ amounts that originate from quantum confinement rather than octahedral tilting. The influence of substituting the monovalent cation depends on the crystalline symmetry of the system [67]. When replacing the original cation in the cubic structure will result in change of lattice parameters as an expansion or as a contraction of the lattice. However, if the structure is tetragonal or orthorhombic, A-cation substitution produces two competing effects regarding both the crystal size and tilting angles of the octahedra. On the one hand, in the case of small A-cation the lattice can shrink, thus strengthening the antibonding overlap between X-p and Pb-s orbitals. On the other hand, small A-cation increases the Pb-X-Pb angle, thus weakening the p-s overlap. It was simulated by Meloni et al., who concluded that the second effect dominates, overall lowering the conduction band and increasing the bandgap [68]. Therefore, it can be concluded that both effects occur. However, since the Bohr diameter of the NCs (which is 5 and 7 nm for CsPbCl3 and CsPbBr3, respectively) [19] is much smaller than the size of NCs in this work, no quantum confinement can take place. Therefore, the optical change is associated with the octahedral tilting. In addition, a control experiment was performed in order to ensure that the addition of Rb+ causes the changes in the optical properties. Cs0.2PbBr3, Cs0.4PbBr3, and Cs0.8PbBr3 without Rb+ were synthesized (importantly, these chemical formulas indicate the composition at the preparation of the solution,

35

2.4 Near-ultraviolet to mid-visible bandgap tuning of mixed-cation

which emphasizes that no Rb+ was used), and their absorption spectra were compared with the corresponding absorption of the syntheses with Rb+ (i.e., Cs0.2Rb0.8PbBr3, Cs0.4Rb0.6PbBr3, and Cs0.8Rb0.2PbBr3 as shown in Figure 23). Overall, it is shown that when Rb+ is added to the NCs the absorbance spectra are blueshifted more than in the case without Rb+. This is shown in Figure 23(a)–(c). This control experiment further shows that the optical changes are a result of the presence of Rb in the NCs. It can be concluded that the amount of Cs in the synthesis affects mainly the amount of NCs that are formed.

(a) 0.7

2.0 Absorbance (a.u)

0.6 Absorbance (a.u)

(b) 2.5

Cs0.4PbBr3 Cs0.4Rb0.6PbBr3

0.5 0.4 0.3 0.2

475

500 525 wavelength (nm)

1.0

(d) 1.0

Cs0.2PbBr3 Cs0.2Rb0.8PbBr3

1.0

0.5

475

500 525 wavelength (nm)

0.0 450

550

Norm. absorbance (a.u)

Absorbance (a.u)

(c) 1.5

0.0 450

1.5

0.5

0.1 0.0 450

Cs0.8PbBr3 Cs0.8Rb0.2PbBr3

550

0.8

500 wavelength (nm)

550

Cs0.8PbBr3 Cs0.2PbBr3 Cs0.4PbBr3

0.6 0.4 0.2 0.0 450

475 500 525 wavelength (nm)

550

Figure 23: (a) Absorbance of Cs0.2PbBr3 and Cs0.2Rb0.8PbBr3. (b) Absorbance of Cs0.4PbBr3 and Cs0.4Rb0.6PbBr3. (c) Absorbance of Cs0.2PbBr3 and Cs0.2Rb0.8PbBr3. (d) Absorbance of Cs0.2PbBr3, Cs0.4PbBr3, and Cs0.8PbBr3 [60].

Moreover, the shape of NCs is alerting when adding more Rb+. It can be related to the change in the crystal phase that may occur due to serious deformations of the perovskite phase into a mix of Cs+- and Rb+-based phases, as the PXRD confirms. Moreover, black dots appear in all the images. A recent work investigated this issue thoroughly and showed that the black dots are Pb0 seeds that nucleate prior

36

2 All-inorganic perovskite nanostructures

to the reaction in the PbX2 flask. According to these findings, the use of Rb+ in the same molar ratio as its Cs+ counterpart is probably insufficient for a complete crystallization; thus more Pb0 seeds appear at higher Rb+ concentrations [68].

2.4.1 Summary This section describes a continuous work of the previous section, where mixedinorganic cations were used at the same NCs. We demonstrate the introduction of Rb+ cation into CsPbX3 NCs. The addition of small Rb+ cation with increasing ratios affects the levels of structural pressure on the inorganic CsPbX3 (X = Cl, Br) perovskite NCs. We provide a detailed characterization of A-cation modifications optically and structurally in the nanoscale, which opens a window for high structural flexibility. The obtained NCs were characterized and found to exhibit high PL-QYs, which are comparable to those of the original CsPbX3 NPs. The optical properties were alerted due to the introduction of the Rb+ cation over the Cs+ due to the octahedral tilting, which affects the antibonding overlap of the Pb2+ and the X− orbitals. TEM images showed square-shaped NCs, and control syntheses without Rb show a variation in optical properties, which further prove the presence of Rb in the NCs. Although calculations of TFs of these NCs predicted that high contents of Rb+ yield are close to the lower limit of the perovskite formability range that can lead to unstable perovskites, the results proved that mixed-cation perovskite NPs are indeed formed, possessing properties that are very similar to the known CsPbX3 NPs. The option to fine-tune the optical properties by the Acation in the nanoscale is one step forward to full understanding of perovskite nanostructures and for future optoelectronic applications.

3 Hybrid perovskite nanostructures This section discusses hybrid perovskite nanostructures, without including inorganic cations. The synthesis of these hybrid perovskite nanostructures is different than the all-inorganic perovskite synthesis. Moreover, these hybrid nanostructures demonstrate various properties and different stability behavior.

3.1 Two-dimensional hybrid perovskite nanorods This section describes the synthesis and properties of 2D OIP nanorods (NRs). The synthesis of the 2D hybrid perovskite NRs is done at low temperature of ca. 80 °C, and the chemical composition of perovskite NRs is based on the formula (C8H17NH3)2(CH3NH3)2Pb3(IxBr1–x)10, 0 > x > 1. The NRs were characterized by XRD, ED, and FFT analysis. TEM micrographs of the synthesized 2D perovskite NRs of the structure R2(MA)n−1MnX3n+1 (where R = OA = octylammonium = C8H17NH3+; MA = CH3NH3+; M = Pb2+; X = I− or Br−, 0 < n < 3) [25, 69, 70] are presented in Figure 24. The figure shows the variation of halide composition in the NRs, and the halide composition did not affect the shape and size of NRs; the average size of the NRs is 2.25 nm ± 0.3 nm width and 11.36 nm ± 2.4 nm length. The inset of Figure 24 shows the FFT of these NRs, and analyzing the FFT reveals the corresponding Miller indices for various compositions. Further analysis of the structure of NRs by XRD and ED shows that the perovskite NRs have 2D perovskite structure combined with 3D tetragonal structure (described in more details below). The introduction of Br into the structure of NRs results in lattice changes from tetragonal lattice parameters of a = b = 8.856 Å , c = 12.674 Å for (OA)2(MA)2Pb3I10 to a = b = 8.611 Å, c = 12.234 Å for (OA)2(MA)2Pb3(IxBr1–x)10(I > Br), a = b = 8.484 Å, c = 12.294 Å for (OA)2(MA)2Pb3(IxBr1–x)10 (I < Br), and a = b = 8.474 Å, c = 11.943 Å for pure (OA)2(MA)2Pb3Br10. The introduction of Br into the perovskite structure was also described previously for bulk perovskite [71], and the changes in the lattice parameters are a result of the smaller ionic radius of the Br− compared with I−. Looking on the XRD pattern of these NRs, we can recognize additional features below 14°. The peak at 14° is the main characteristic of the 3D perovskite structure, while the peaks below 14° are characteristic of 2D perovskite structure, and they are originated from the diffraction of the X-rays with the (00n) facets of the 2D perovskite crystal (n = 2, 4, 6, . . .) as indicated by Mitzi et al. [73]. We used the (002) peak (the angles range between 5.3° and 6.5° based on the concentration of I to Br) to calculate the d-spacing (dð002Þ ) and the lattice https://doi.org/10.1515/9783110602173-004

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3 Hybrid perovskite nanostructures

Figure 24: TEM images and inset show FFTs of the NRs made of various halide compositions [72].

parameter “c” of the resulting perovskite. We found that the lattice parameter “c” of the various NR composition is larger than the tetragonal “c” (as an example, for the pure-iodide sample, the ctetragonal is 12.674 Å while for the c2D it is 27.19 Å; thus, c2D ≈ 2ctetragonal). This is a confirmation for the 2D nature of the NRs. Absorbance and PL measurements were performed on the 2D hybrid perovskite NRs (see Figure 25(b) and (c)). It can be seen from the absorbance spectra that the full iodide and full bromide NRs, that is, (OA)2(MA)2Pb3I10 NRs and (OA)2(MA)2Pb3Br10 NRs, absorb only until ⁓650 and ⁓530 nm, respectively (Figure 25(b)), which is not as the absorbance of the bulk of MAPbI3 and MAPbBr3 [2]. The difference in the absorbance is attributed to two main contributions. The first is related to the use of octylammonium as ligand in the synthesis of NRs. It can suggest that the octylammonium is not incorporate into the NRs perovskite structure due to the reason that it is larger than the MA. The crystal growth is limited by the attachment of the octylammonium through its alkyl chain to the NRs surface [74]. As a result of this the perovskite growth is formed in a 2D perovskite structure, which causes a shift in the absorbance and PL to shorter wavelength (higher energies) and hence to larger bandgaps. The absorbance matches the reported absorbance of layered perovskite, where n = 3 [16, 23, 75]. The second reason for the shift in the optical properties is related to the halide exchange, which includes the introduction of bromide ions into the (OA)2(MA)2Pb3I10 NRs (and the exclusion of iodide ions, respectively). The orbitals of Br(4p) together with the Pb(6s) orbitals determine the energy of the absorption peak. The PbI2 conduction band mainly contains Pb(6p) orbitals, where its VB is composed of Pb(6s) orbitals and I(5p) orbitals. The PbI2 transitions are similar to the transitions in MAPbIxBr3–x (0 > x > 3) [76, 77]. Now, the

3.1 Two-dimensional hybrid perovskite nanorods

(a)

(b)

(c)

(d)

39

Figure 25: (a) Photo of the emitting NR dispersions composed of various halide mixtures. (b) Absorbance spectra of the 2D perovskite NRs made of various halide compositions. (c) PL spectra of the 2D perovskite NRs made of various halide compositions. (d) PL-QY and bandgap energies measured and calculated by the Tauc plot of the perovskite NR compositions [72].

energy of the Br(4p) orbitals is lower than the energy of Pb(6s) orbitals; therefore, the absorption peak position of (OA)2(MA)2Pb3(IxBr1–x)10, 0 > x > 1, is influenced and shifted to higher energy [78]. Figure 25(b) and (c) shows a shift in the absorbance and PL to higher energies. Figure 25(a) shows an image with the emission of NRs having different compositions from iodide to bromide. These emission colors agree well with the PL measurement. The bandgap of various NR compositions was calculated using Tauc plot measurements. The Eg values are presented in Figure 25(d) (red). In the case of the bulk MAPbX3 perovskites, as the Br/I ratio increases the materials’ Eg also rises. The range of bandgap of these NRs is between 1.9 and 2.26 eV due to the dimensional changes. Figure 25(d) (black) shows the PL-QY of various NR compositions. The (OA)2(MA)2Pb3I10 NRs and the (OA)2(MA)2Pb3(IxBr1–x)10 (I > Br) NRs show a PL-QY

40

3 Hybrid perovskite nanostructures

close to 30% that is much higher than the corresponding PL-QY of MAPbI3 bulk, which is close to 2–5%. Feldmann et al. [47] have reported low PL-QY for bromide-based nanostructures similar to the observation in this study. When the Br concentration is increased, the PL-QY is decreased.

3.1.1 Role of the ligands The main factor influencing the formation of NRs is the organic moieties present in the synthesis. Three organic molecules used in the synthesis are MA iodide/bromide (MAI/MABr), octylammonium iodide (OAI), and OAc. As described earlier, the long chain of the OAI cannot be incorporated into the perovskite crystal. Therefore, it is attached to specific sites on the perovskite surface inhibiting the growth in a particular direction. The OAc also plays an important role in the formation of these 2D NRs, which is discussed later. In order to study the effect of the ligands on the formation of the 2D NRs, we used several ratios of OAc to OAI. In this chapter, NRs showed a ratio of OAI/OAc = 0.186. Therefore, ratios of 100% OAc, OAI/OAc = 0.075 (lower than the standard ratio), OAI/OAc = 0.250 (higher than the standard ratio) and 100% OAI were studied. Figure 26(a)–(e) shows TEM images for the different ligand ratios. In the case of 100% OAc (without OAI in the solution), there was no indication for the formation of nanostructures; the perovskite was formed as a bulk that precipitated at the bottom of the vial (Figure 26(a) shows empty TEM grid). Figure 26(d) shows a ratio of OAI/OAc = 0.075, where the formation of QDs was observed, although their concentration was low. In the case of higher ratios, for example, OAI/OAc = 0.250, both QDs and NRs were observed, which is indicated in Figure 26(d). As for the case of just OAI (100%), a high concentration of QDs was observed (Figure 26(e)). Several conclusions can be summarized from this ligand study. At 100% OAc, no QDs were observed, which means that they cannot stabilize the QDs alone; on the other hand, in the case of 100% OAI, QDs were formed which means that OAI is enough to stabilize their growth. The presence of iodide in the OAI assists in the attachment of this ligand to the QDs/NRs surface. This kind of attachment is called “chemisorption,” as the interaction between the ligand and the surface leads to the formation of chemical interaction. In the case of OAc, the interaction is physisorption, which means that there is just physical adsorption between the surface of the particles and the OAc. Looking on to other ligand ratios, that is, ratio of 0.250 (excess of OAI in relation to the standard) and ratio of 0.075 (excess of OAc in relation to the standard). In the first case, QDs and NRs are formed, indicated by the two

(f)

(e)

(c)

(d)

Figure 26: (a)–(f) TEM images of MAPbI3 QDs/NRs were prepared at various OAI/OAc ratios: (a) 100% OAc; (b) OAI/OAc = 0.075; (c) OAI/OAc = 0.186; (d) OAI/OAc = 0.250; the two figures are taken from the same sample in different areas; (e) 100% OAI. (f) Schematic illustration of the suggested formation mechanism of NRs; the OAI and the OAc are presented schematically in the figure, while the black dots represent methylammonium and the brown rhombus represent PbI2 [72].

(b)

(a)

3.1 Two-dimensional hybrid perovskite nanorods

41

42

3 Hybrid perovskite nanostructures

images in Figure 26(d) which were taken from the same sample. For a ratio of 0.075, the formation of NRs was rare. It can be concluded that the OAI is attached to the surface of NRs, while the OAc has an important role by shaping these QDs into NRs. As described earlier, the OAI attachment is much stronger to the surface of NRs than that of the OAc since it fills the octahedral hole. Figure 26(f) shows a schematic illustration of the ligand’s role in the synthesis of NRs. It is clear that the presence of the two ligands is essential for the formation of NRs. The OAI has a high affinity to the perovskite surface more than the OAc, which absorbs just physically to the surface of NRs. Therefore, pure NRs were formed just in a specific ratio of OAI/OAc = 0.186, which is enough to stabilize the structure and to orient the growth of NRs at the same time.

3.1.2 Summary In this section, we present the synthesis of 2D hybrid perovskite NRs of the structure (C8H17NH3)2(CH3NH3)2Pb3(IxBr1–x)10, 0 > x > 1. The synthesis of these hybrid perovskite nanostructure is different than the Cs-based nanostructures. The synthesis is done at low temperature with the use of antisolvent. We showed that the optical properties can be controlled by varying the halides, in addition a very good size distribution and strong PL were observed. It was found that these NRs exhibit 2D structure due to the use of OAI as the ligand. The wider bandgap observed for these NRs is related to the characteristics of their 2D structure, due to the use of OAI in their synthesis. XRD, ED, and FFT provide evidence for the crystallographic structure of these NRs. Finally, we investigate the role of ligands in the formation of these NRs. It was found that OAI is essential for the formation of NRs, where a specific ratio of OAI/OAc provides the shape and the size of the NRs.

3.2 The effect of the alkylammonium ligands length on OIP NPs In this section, we study how the length of the alkylammonium cation influences the optical and physical properties of OIP NPs. In this study, three alkylammonium cations (octyl, dodecyl, and octadecyl ammonium, C8, C12, and C18, respectively) that function as the ligands were investigated. The properties of the NPs were compared to the properties of 2D perovskite films using the same alkylammonium cations. In the case of NPs, these alkylammonium cations function as ligands while in the case of 2D perovskite thin films, these function as barrier molecules.

3.2 The effect of the alkylammonium ligands length on OIP NPs

43

The ligands used in this study for the OIP NPs are shown in Figure 27(a). Both bromide- and iodide-based NPs were synthesized using each of the three linear alkylammonium cations. (a)

(b)

Figure 27: (a) The alkylammonium ligands used in this work are octylammonium (C8), dodecylammonium (C12), and octadecylammonium (C18). The gray spheres represent carbon, the blue spheres nitrogen, and the white spheres hydrogen. (b) The size of NPs as a function of the alkylammonium length for both bromide-based NPs (green) and iodide-based NPs (brown). The average edge size of the C8-, C12-, and C18-based NPs was 3.3 ± 0.5, 8 ± 2, and 11 ± 4 nm for the bromide-based NPs, and 2.2 ± 0.4 (width), 5 ± 2, and 13 ± 4 nm for the iodide-based NPs, respectively [79].

The syntheses of the OIP NPs were carried out under ambient atmosphere based on our previous report [72], including the following steps: (1) injection of lead halide (PbI2 or PbBr2) and methylammonium halide (MAI or MABr) solutions in DMF into a hot medium (80 °C) that contained ODE, OAc, and alkylammonium halide (C8-I, C8-Br, C12-I, C12-Br, or C18-I, C18-Br) under vigorous stirring. (2) Addition of chloroform triggered the formation of NPs. (3) Centrifugation of dispersions led to sedimentation of bulky by-products and the acquisition of clear dispersions. Here, the alkylammonium cations are defined as “ligands”; however, the meaning of this term is different from the usual meaning in the colloidal synthesis of NPs. In this case, the interactions of alkylammonium with the surface of NPs and the interactions of MA (inside the core of the NP) with the inorganic octahedra are ionic. Both the MA and alkylammonium have hydrogen bonds with the halogens that surround the metal cation. In general, common ligands stabilize the NPs due to surface orbitals of the atoms by chemical or physical bonds. Therefore, in contrast to common ligands the alkylammonium cations affect the electronic structure of NPs [80]. Figure 28 presents the TEM images of the synthesized NPs using different alkylammonium cations. It is shown that in all cases a clear and defined NPs

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3 Hybrid perovskite nanostructures

(a)

(b)

(c)

(d)

(e)

(f)

Figure 28: TEM images of OIP NPs that were synthesized using different lengths of alkylammonium with bromide or iodide. The alkylammonium cations in each figure are the following: (a) C8-Br, (b) C12-Br, (c) C18-Br, (d) C8-I, (e) C12-I, and (f) C18-I [79].

(nanocubes or NRs) were observed in the synthesis of OIP NPs which is contrary to all-inorganic perovskite NPs. This is trivial to get well-defined NPs under ambient atmosphere. There are several reports in the literature, which show OIP NPs prepared under ambient conditions, where their shape and structure were imperfect which did not assemble into the perovskite’s typical mosaic-like array [81, 82]. The stability of NPs to the TEM electron beam is observed in Figure 28. The OIP NPs are assembled in arrays, which are typical for inorganic perovskite NPs that have cubic or sheet-like shapes, as well as a mosaic-like assembly [16, 82]. The length of the alkylammonium was found to affect the size of NPs, as shown in Figures 27(b) and 28(a)–(f). The correlation is clear, as the length of the alkylammonium increased and the average size of the NPs increased. Another interesting observation is that when using iodie-C8 versus bromide-C8 different morpholgy was observed (i.e. NRs Vs. nanocubes), this may be result from different surface energy [74]. Importantly, short alkyl moieties undergo weaker van der Waals (VDW) interactions than the long alkylic moieties. The tails of the alkylammonium ligands prefer to increase VDW interactions among themselves. The surface energy of the NPs changes with the ligand length, that is, γC18 > γC12 > γC8 . Therefore, when the surface energy is high, the NP surface-to-volume ratio is low. This explains why a long alkylammonium forms larger NPs than does a short alkylammonium.

3.2 The effect of the alkylammonium ligands length on OIP NPs

45

The absorption and PL spectra for the bromide- and iodide-based NPs with different alkylammonium lengths are shown in Figure 29(a) and (d), and (b) and (e), respectively, while the emission of NPs under ultraviolet (UV) light (λ = 254 nmÞ is observed in Figure 29(c) and (f). The difference in the optical properties between bromide and iodide NPs is due to the different p orbitals of the halides (4p of bromide vs. 5p of iodide) as discussed in the previous sections. In addition to that the absorption onset and the PL maximum were shifted to higher energies than those of the corresponding bulk material [2] due to the confinement effect that takes place in these NPs. The alkylammonium ligands strongly affect the electronic structure of the OIP NPs due to the strong ionic interaction between the ammonium group and the inorganic octahedra. The cation–halogen interactions increase where the metal–halide interactions are suppressed in the transition of the perovskite to low dimensions. Importantly, the difference between the dielectric constants of the inorganic part and the organic part (alkylammonium) is predominant; thus, the bandgap energy of the perovskite increases. Moreover, the optical properties (absorption onset and the PL maximum) varied by varying the alkylammonium length, which can be due to several reasons: (1) the variation of the ligands’ dielectric constant; (2) the angle between the ligand and the perovskite surface, which defines the level of the crystalline structure’s distortion; and (3) the dominance of VDW interactions between adjacent alkylammonium ligands at the surface of NPs. The bonding angle is independent of the alkylammonium length as was reported [83], which role out the second reason. Moreover, the dielectric constants of C8, C12, and C18 are similar; thus, this factor also cannot explain the significant shift that was observed in the optical measurements [84]. Based on this, the main factor that affects the optical properties of NPs is the different strength of VDW interactions among the alkylammonium ligands of various lengths.

3.2.1 Study of the VDW interactions A comparison is done between NPs with a difference that films of 2D perovskite having the formula (RNH3)2(MA)n–1PbnX3n+1 (where n = 1, 2, or 3, R is an aliphatic chain of lengths C8, C12, or C18, and X is bromide or iodide) were fabricated. The absorption spectra of 2D films present various number of absorption peaks in each spectrum, and distinctions in peak dominance (despite the identical n-value, not shown). These results show that the optical properties

(e)

(d)

(f)

(c)

Figure 29: Absorption spectra of (a) bromide- and (d) iodide-based OIP NPs were synthesized using different lengths of alkyl ammonium ligands. PL spectra of (b) bromide- and (e) iodide-based OIP NPs were synthesized using different lengths of alkyl ammonium ligands. Pictures of (c) bromide- and (f) iodide-based OIP NP dispersions (in chloroform) were synthesized using different lengths of alkylammonium ligands under UV light and under white light (inset) [79].

(b)

(a)

46 3 Hybrid perovskite nanostructures

3.2 The effect of the alkylammonium ligands length on OIP NPs

47

change among similar n-values only due to the change in the lengths of alkylammonium. The alkylammonium cation affects the assembly process of the perovskite layers due to their VDW interactions. Figure 30 presents the pXRD patterns of 2D perovskite films. Theoretically, the d001 value expands as the length of alkylammonium increases. Thus, the angle of the (001) diffraction peak is shifted toward low angles [84]. Indeed, this dependence was observed, as shown in Figure 30(a)–(f). In addition, an increase in n value is also a reason for the shift in the (001) peak [85]. In the case of n = 1 and n = 3 layers of C8-Br, their peak appeared at different angles (Figure 30(a)). The first peak of the n = 1 layer appeared at 4.2 , which indicates that d001 = 21.02 Å, whereas n = 3 layer exhibited two peaks, the first at 2.7° (d001 = 32.2 Å) and the second at 3.4 (d001 = 26.26 Å). Importantly, the layers of n > 1 exhibited weak 3D perovskite peaks that appear at 14–15 and 28–29 , since the periodic inorganic parts also reflect the X-ray photons [85]. These 3D perovskite peaks can be observed in the diffractograms of the n = 3 2D perovskites (see the red diffractograms in Figure 30(a)–(f)). A single 2D perovskite crystalline structure can be observed in the diffractogram, where a single set of (00n) peaks appeared. This set of peaks appears due to the reflection of X-rays from the multiplicities of (001) plane by integers. However, as shown in Figure 30(a)–(f) instead of a single set of peaks, a few sets of peaks appeared for n > 1 (see the few peaks at angles below 5° in the red curves in Figure 30(a)–(f)). This observation is valid for both halides (Br− and I−) and for all lengths of alkylammonium (C8, C12, and C18). These peaks are more pronounced in the short alkylammonium than in the long alkylammonium. An important conclusion is related to the formation of n = 1, which tends to be formed using the long alkylammonium cations, whereas the formation of n > 1 perovskite structures becomes undesirable [28], as shown in Figure 30(a)–(f), which shows that the similarity between the diffractograms of n = 1 and n = 3 perovskite films increased as the alkylammonium ligand elongated. These results show that long alkylammonium cations direct the crystalline organization toward lower n-values. As the length of the alkylammonium increases, the VDW interaction among adjacent hydrocarbons enhances. This is in good agreement with the TEM results. These observations help us to understand the absorption onset of small NPs, which is shifted toward shorter wavelengths. Short ligands form NPs that have large surface-to-volume ratios due to the case that VDW interactions among adjacent ligands are moderate. Accordingly, the numerous alkylammonium–halide interactions strongly affect the electronic structure, and lead to bandgap widening.

(e)

(d)

(f)

(c)

Figure 30: Powder XRD diffractograms of 2D perovskite layers of the formula (RNH3)2(MA)n–1PbnX3n+1, where R = C8, C12, or C18, n = 1 (black) or 3 (red) and X = Br− or I− [79].

(b)

(a)

48 3 Hybrid perovskite nanostructures

3.2 The effect of the alkylammonium ligands length on OIP NPs

49

3.2.2 Summary In this section, we demonstrated the synthesis of OIP NPs using different alkylammonium lengths. We studied the effect of the alkylammonium length on the optical, physical, and structural properties. It was concluded that the VDW interactions had a strong effect on the formation, size, and optical properties of NPs. In order to elucidate the role of VDW interactions on the NPs we compared thin film versus NPs using the same alkylammonium ligands as the building blocks. This study helps us to understand the role of the alkylammonium length on the properties and morphology of NPs. The VDW interactions among the different alkylammonium ligands are the main factor for the differences in the optical and physical properties of the synthesized OIP NPs. The PL-QY of the NPs varied from 30% to 60%, depending on the halide and the alkylammonium ligand length. Understanding how the length of the alkylammonium affects the properties of the NPs is essential for future implementation of these NPs in optoelectronic applications.

4 Summary and outlook This book concentrated on all-inorganic and hybrid perovskite nanostructures. It described the difference in synthesis of these nanostructures, how the shape and size can be controlled, and how the chemical manipulation can be performed. The first section shows the synthesis of Cs-based NPs with a study on their kinetic growth. Following that we showed the possibility to synthesize different shapes of perovskite nanostructures, from long NWs to NRs with the ability to tune their optical properties by halide exchange during the synthesis. The second section summarized by demonstrating the possibility to use a smaller inorganic cation (i.e., Rb+) in the NPs, which does not form a perovskite structure in the bulk. By choosing the smallest halide (chloride in this case), we were able to show Rubidium lead chloride NPs with UV absorption. The next stage was to tune the optical properties of these NPs to the visible and nearinfrared regions. Therefore, we mixed Rb+ and Cs+ together in the same NPs. By varying the composition, we were able to tune the optical properties from blue to the visible regions. The most important advantage in this case is the fine-tuning of the optical properties that cannot be achieved by exchanging the halide. The change in the size of the inorganic cation influenced the Pb–X–Pb angle. When this angle is more distorted (whether it is larger or smaller than the ideal 180° angle) and the octahedral tilting is larger, the bandgap will be changed. The third section concentrated on hybrid nanostructures, discussing their synthesis. First, we showed the synthesis and characterization of NR-based perovskite with several unit cell thickness and extremely good size distribution. The ligands in the hybrid perovskite nanostructures played a major role influencing the shape and the stability of these NPs. We showed a correlation between 2D perovskite films to the hybrid perovskite NPs using the same alkylammonium cations as ligands. It was concluded that the VDW interactions controlled the size and shape of NPs. To sum up, the semiconductor OIP opens a fascinating field in the nanoscale materials. This area is not similar to the well-known semiconductor nanomaterials, where their synthesis and optical perovskite were due to the quantum size effect. Here the chemical composition controls the properties. Moreover, the synthesis of these nanostructures is easier and simple. The perovskite nanostructures have high PL-QY, which opens the way to a variety of highly efficient optoelectronic applications. On the other hand, there are numerous fundamental properties that have not been investigated so far. This positions these all-inorganic and hybrid perovskite nanostructures in the forefront of science. https://doi.org/10.1515/9783110602173-005

Acknowledgments The author thanks all the students in his research group along the last years, especially Daniel Amgar, Sigalit Aharon, Miri Koolik, Tal Binyamin, Stav Rahmany, and Dr Malgorzata Wiezbrowska, for their involvement in experiments and analyzing the results and discussions until achieving these wonderful results. This work was financially supported by the Israel Ministry of Science and the Israel Science Foundation.

https://doi.org/10.1515/9783110602173-006

References [1]

Stoumpos, C. C.; Malliakas, C. D.; Kanatzidis, M. G. Semiconducting Tin and Lead Iodide Perovskites with Organic Cations: Phase Transitions, High Mobilities, and near-Infrared Photoluminescent Properties. Inorg. Chem. 2013, 52(15), 9019–9038. https://doi.org/10.1021/ic401215x. [2] Noh, J. H.; Im, S. H.; Heo, J. H.; Mandal, T. N.; Seok, S. Il. Chemical Management for Colorful, Efficient, and Stable Inorganic–Organic Hybrid Nanostructured Solar Cells. Nano Lett. 2013, 13(4), 1764–1769. https://doi.org/10.1021/nl400349b. [3] Hao, F.; Stoumpos, C. C.; Chang, R. P. H.; Kanatzidis, M. G. Anomalous Band Gap Behavior in Mixed Sn and Pb Perovskites Enables Broadening of Absorption Spectrum in Solar Cells. J. Am. Chem. Soc. 2014, 136(22), 8094–8099. https://doi.org/10.1021/ja5033259. [4] Filip, M. R.; Eperon, G. E.; Snaith, H. J.; Giustino, F. Steric Engineering of Metal-Halide Perovskites with Tunable Optical Band Gaps. Nat. Commun. 2014, 5(1), 5757. https://doi.org/10.1038/ncomms6757. [5] Amat, A.; Mosconi, E.; Ronca, E.; Quarti, C.; Umari, P.; Nazeeruddin, M. K.; Grätzel, M.; De Angelis, F. Cation-Induced Band-Gap Tuning in Organohalide Perovskites: Interplay of Spin–Orbit Coupling and Octahedra Tilting. Nano Lett. 2014, 14(6), 3608–3616. https://doi.org/10.1021/nl5012992. [6] Stoumpos, C. C.; Frazer, L.; Clark, D. J.; Kim, Y. S.; Rhim, S. H.; Freeman, A. J.; Ketterson, J. B.; Jang, J. I.; Kanatzidis, M. G. Hybrid Germanium Iodide Perovskite Semiconductors: Active Lone Pairs, Structural Distortions, Direct and Indirect Energy Gaps, and Strong Nonlinear Optical Properties. J. Am. Chem. Soc. 2015, 137(21), 6804–6819. https://doi.org/10.1021/jacs.5b01025. [7] Liu, W.; Lin, Q.; Li, H.; Wu, K.; Robel, I.; Pietryga, J. M.; Klimov, V. I. Mn2+ -Doped Lead Halide Perovskite Nanocrystals with Dual-Color Emission Controlled by Halide Content. J. Am. Chem. Soc. 2016, 138(45), 14954–14961. https://doi.org/10.1021/jacs.6b08085. [8] Parobek, D.; Roman, B. J.; Dong, Y.; Jin, H.; Lee, E.; Sheldon, M.; Son, D. H. Exciton-toDopant Energy Transfer in Mn-Doped Cesium Lead Halide Perovskite Nanocrystals. Nano Lett. 2016, 16(12), 7376–7380. https://doi.org/10.1021/acs.nanolett.6b02772. [9] Liu, H.; Wu, Z.; Shao, J.; Yao, D.; Gao, H.; Liu, Y.; Yu, W.; Zhang, H.; Yang, B. CsPbxMn1–xCl3 Perovskite Quantum Dots with High Mn Substitution Ratio. ACS Nano 2017, 11(2), 2239–2247. https://doi.org/10.1021/acsnano.6b08747. [10] Duong, T.; Mulmudi, H. K.; Shen, H.; Wu, Y.; Barugkin, C.; Mayon, Y. O.; Nguyen, H. T.; Macdonald, D.; Peng, J.; Lockrey, M.; et al. Structural Engineering Using Rubidium Iodide as a Dopant under Excess Lead Iodide Conditions for High Efficiency and Stable Perovskites. Nano Energy. 2016, 30, 330–340. https://doi.org/10.1016/J.NANOEN.2016.10.027. [11] Duong, T.; Wu, Y.; Shen, H.; Peng, J.; Fu, X.; Jacobs, D.; Wang, E.-C.; Kho, T. C.; Fong, K. C.; Stocks, M.; et al. Rubidium Multication Perovskite with Optimized Bandgap for Perovskite-Silicon Tandem with over 26% Efficiency. Adv. Energy Mater. 2017, 7(14), 1700228. https://doi.org/10.1002/aenm.201700228. [12] Saliba, M.; Matsui, T.; Domanski, K.; Seo, J.-Y.; Ummadisingu, A.; Zakeeruddin, S. M.; Correa-Baena, J.-P.; Tress, W. R.; Abate, A.; Hagfeldt, A.; et al. Incorporation of Rubidium Cations into Perovskite Solar Cells Improves Photovoltaic Performance. Science 2016, 354(6309), 206–209. https://doi.org/10.1126/science.aah5557.

https://doi.org/10.1515/9783110602173-007

54

References

[13] Linaburg, M. R.; McClure, E. T.; Majher, J. D.; Woodward, P. M. Cs1– xRbxPbCl3 and Cs1– xRbxPbBr3 Solid Solutions: Understanding Octahedral Tilting in Lead Halide Perovskites. Chem. Mater. 2017, 29(8), 3507–3514. https://doi.org/10.1021/acs.chemmater.6b05372. [14] https://www.nrel.gov/pv/assets/pdfs/pv-efficiencies-07-17-2018.pdf. [15] Yang, W. S.; Park, B.-W.; Jung, E. H.; Jeon, N. J.; Kim, Y. C.; Lee, D. U.; Shin, S. S.; Seo, J.; Kim, E. K.; Noh, J. H.; et al. Iodide Management in Formamidinium-Lead-Halide-Based Perovskite Layers for Efficient Solar Cells. Science 2017, 356(6345), 1376–1379. https://doi.org/10.1126/science.aan2301. [16] Protesescu, L.; Yakunin, S.; Bodnarchuk, M. I.; Krieg, F.; Caputo, R.; Hendon, C. H.; Yang, R. X.; Walsh, A.; Kovalenko, M. V. Nanocrystals of Cesium Lead Halide Perovskites (CsPbX3, X = Cl, Br, and I): Novel Optoelectronic Materials Showing Bright Emission with Wide Color Gamut. Nano Lett. 2015, 15(6), 3692–3696. https://doi.org/10.1021/nl5048779. [17] Nedelcu, G.; Protesescu, L.; Yakunin, S.; Bodnarchuk, M. I.; Grotevent, M. J.; Kovalenko, M. V. Fast Anion-Exchange in Highly Luminescent Nanocrystals of Cesium Lead Halide Perovskites (CsPbX 3, X = Cl, Br, I). Nano Lett. 2015, 15(8), 5635–5640. https://doi.org/10.1021/acs.nanolett.5b02404. [18] Akkerman, Q. A.; D’Innocenzo, V.; Accornero, S.; Scarpellini, A.; Petrozza, A.; Prato, M.; Manna, L. Tuning the Optical Properties of Cesium Lead Halide Perovskite Nanocrystals by Anion Exchange Reactions. J. Am. Chem. Soc. 2015, 137(32), 10276–10281. https://doi.org/10.1021/jacs.5b05602. [19] Protesescu, L.; Yakunin, S.; Kumar, S.; Bär, J.; Bertolotti, F.; Masciocchi, N.; Guagliardi, A.; Grotevent, M.; Shorubalko, I.; Bodnarchuk, M. I.; et al. Dismantling the “Red Wall” of Colloidal Perovskites: Highly Luminescent Formamidinium and Formamidinium–Cesium Lead Iodide Nanocrystals. ACS Nano 2017, 11(3), 3119–3134. https://doi.org/10.1021/acsnano.7b00116. [20] Goldschmidt, V. M. Die Gesetze Der Krystallochemie. Naturwissenschaften. 1926, 14(21), 477–485. https://doi.org/10.1007/BF01507527. [21] Kieslich, G.; Sun, S.; Cheetham, A. K. An Extended Tolerance Factor Approach for Organic-Inorganic Perovskites. Chem. Sci. 2015, 6(6), 3430–3433. https://doi.org/10.1039/c5sc00961h. [22] Travis, W.; Glover, E. N. K.; Bronstein, H.; Scanlon, D. O.; Palgrave, R. G. On the Application of the Tolerance Factor to Inorganic and Hybrid Halide Perovskites: A Revised System. Chem. Sci. 2016, 7(7), 4548–4556. https://doi.org/10.1039/c5sc04845a. [23] Green, M. A.; Ho-Baillie, A.; Snaith, H. J. The Emergence of Perovskite Solar Cells. Nat. Photonics. Nature Publishing Group July 2014, pp 506–514. https://doi.org/10.1038/nphoton.2014.134. [24] Bekenstein, Y.; Koscher, B. A.; Eaton, S. W.; Yang, P.; Alivisatos, A. P. Highly Luminescent Colloidal Nanoplates of Perovskite Cesium Lead Halide and Their Oriented Assemblies. J. Am. Chem. Soc. 2015, 137(51), 16008–16011. https://doi.org/10.1021/ jacs.5b11199. [25] Akkerman, Q. A.; Motti, S. G.; Srimath Kandada, A. R.; Mosconi, E.; D’Innocenzo, V.; Bertoni, G.; Marras, S.; Kamino, B. A.; Miranda, L.; De Angelis, F.; et al. Solution Synthesis Approach to Colloidal Cesium Lead Halide Perovskite Nanoplatelets with Monolayer-Level Thickness Control. J. Am. Chem. Soc. 2016, 138(3), 1010–1016. https://doi.org/10.1021/jacs.5b12124.

References

55

[26] Vybornyi, O.; Yakunin, S.; Kovalenko, M. V. Polar-Solvent-Free Colloidal Synthesis of Highly Luminescent Alkylammonium Lead Halide Perovskite Nanocrystals. Nanoscale. 2016, 8(12), 6278–6283. https://doi.org/10.1039/C5NR06890H. [27] Protesescu, L.; Yakunin, S.; Bodnarchuk, M. I.; Bertolotti, F.; Masciocchi, N.; Guagliardi, A.; Kovalenko, M. V. Monodisperse Formamidinium Lead Bromide Nanocrystals with Bright and Stable Green Photoluminescence. J. Am. Chem. Soc. 2016, 138(43), 14202–14205. https://doi.org/10.1021/jacs.6b08900. [28] Amgar, D.; Aharon, S.; Etgar, L. Inorganic and Hybrid Organo-Metal Perovskite Nanostructures: Synthesis, Properties, and Applications. Adv. Funct. Mater. 2016, 26(47), 8576–8593. https://doi.org/10.1002/adfm.201603752. [29] Saparov, B.; Mitzi, D. B. Organic–Inorganic Perovskites: Structural Versatility for Functional Materials Design. Chem. Rev. 2016, 116(7), 4558–4596. https://doi.org/10.1021/acs.chemrev.5b00715. [30] Mitzi, D. B. Templating and Structural Engineering in Organic–Inorganic Perovskites. J. Chem. Soc. Dalt. Trans. 2001, 0(1), 1–12. https://doi.org/10.1039/b007070j. [31] Huang, H.; Zhao, F.; Liu, L.; Zhang, F.; Wu, X.; Shi, L.; Zou, B.; Pei, Q.; Zhong, H. Emulsion Synthesis of Size-Tunable CH 3 NH 3 PbBr 3 Quantum Dots: An Alternative Route toward Efficient Light-Emitting Diodes. ACS Appl. Mater. Interfaces 2015, 7(51), 28128–28133. https://doi.org/10.1021/acsami.5b10373. [32] Boles, M. A.; Ling, D.; Hyeon, T.; Talapin, D. V. The Surface Science of Nanocrystals. Nat. Mater. 2016, 15(2), 141–153. https://doi.org/10.1038/nmat4526. [33] Eperon, G. E.; Stranks, S. D.; Menelaou, C.; Johnston, M. B.; Herz, L. M.; Snaith, H. J. Formamidinium Lead Trihalide: A Broadly Tunable Perovskite for Efficient Planar Heterojunction Solar Cells. Energy Environ. Sci. 2014, 7(3), 982. https://doi.org/10.1039/c3ee43822h. [34] Koolyk, M.; Amgar, D.; Aharon, S.; Etgar, L. Kinetics of Cesium Lead Halide Perovskite Nanoparticle Growth; Focusing and de-Focusing of Size Distribution. Nanoscale 2016, 8(12), 6403–6409. https://doi.org/10.1039/C5NR09127F. [35] Chang, Y. H.; Park, C. H.; Matsuishi, K. First-Principles Study of the Structural and the Electronic Properties of the Lead-Halide-Based Inorganic-Organic Perovskites (CH3NH3) PbX3 and CsPbX3 (X = Cl, Br, I). J. Korean Phys. Soc. 2004, 44(4), 889–893. https://doi.org/10.3938/jkps.44.889. [36] Murtaza, G.; Ahmad, I. First Principle Study of the Structural and Optoelectronic Properties of Cubic Perovskites CsPbM3 (M=Cl, Br, I). Phys. B Condens. Matter 2011, 406(17), 3222–3229. https://doi.org/10.1016/J.PHYSB.2011.05.028. [37] Xiaogang Peng; J. Wickham, and; Alivisatos*, A. P. Kinetics of II-VI and III-V Colloidal Semiconductor Nanocrystal Growth: “Focusing” of Size Distributions. 1998. https://doi.org/10.1021/JA9805425. [38] Matsunaga, N.; Hori, M.; Nagashima, A. Gaseous Diffusion Coefficients of Methyl Bromide and Methyl Iodide into Air, Nitrogen, and Oxygen. Heat Transf. Res. 2009, 38(6), 361–369. https://doi.org/10.1002/htj.20255. [39] Yu, P. Y.; Cardona, M. Fundamentals of Semiconductors; Springer Berlin Heidelberg: Berlin, Heidelberg, 1996. https://doi.org/10.1007/978-3-662-03313-5. [40] Cottingham, P.; Brutchey, R. L. On the Crystal Structure of Colloidally Prepared CsPbBr3 Quantum Dots. Chem. Commun. 2016, 52(30), 5246–5249. https://doi.org/10.1039/C6CC01088A.

56

References

[41] Liu, M.; Johnston, M. B.; Snaith, H. J. Efficient Planar Heterojunction Perovskite Solar Cells by Vapour Deposition. Nature 2013, 501. https://doi.org/10.1038/nature12509. [42] Amgar, D.; Stern, A.; Rotem, D.; Porath, D.; Etgar, L. Tunable Length and Optical Properties of CsPbX3 (X = Cl, Br, I) Nanowires with a Few Unit Cells. Nano Lett. 2017, 17(2), 1007–1013. https://doi.org/10.1021/acs.nanolett.6b04381. [43] Myers, R. T. The Strength of the Hydrohalic Acids. J. Chem. Educ. 1976, 53(1), 17. https://doi.org/10.1021/ed053p17. [44] Schmid, R.; Miah, A. M. The Strength of the Hydrohalic Acids. J. Chem. Educ. 2001, 78 (1), 116. https://doi.org/10.1021/ed078p116. [45] Zhang, D.; Yang, Y.; Bekenstein, Y.; Yu, Y.; Gibson, N. A.; Wong, A. B.; Eaton, S. W.; Kornienko, N.; Kong, Q.; Lai, M.; et al. Synthesis of Composition Tunable and Highly Luminescent Cesium Lead Halide Nanowires through Anion-Exchange Reactions. J. Am. Chem. Soc. 2016, 138(23), 7236–7239. https://doi.org/10.1021/jacs.6b03134. [46] Nikl, M.; Mihokova, E.; Nitsch, K.; Somma, F.; Giampaolo, C.; Pazzi, G.; Fabeni, P.; Zazubovich, S. Photoluminescence of Cs4PbBr6 Crystals and Thin Films. Chem. Phys. Lett. 1999, 306 (5–6), 280–284. https://doi.org/10.1016/S0009-2614(99)00477-7. [47] Sichert, J. A.; Tong, Y.; Mutz, N.; Vollmer, M.; Fischer, S.; Milowska, K. Z.; García Cortadella, R.; Nickel, B.; Cardenas-Daw, C.; Stolarczyk, J. K.; et al. Quantum Size Effect in Organometal Halide Perovskite Nanoplatelets. Nano Lett. 2015, 15(10), 6521–6527. https://doi.org/10.1021/acs.nanolett.5b02985. [48] Yuan, Z.; Shu, Y.; Tian, Y.; Xin, Y.; Ma, B. A Facile One-Pot Synthesis of Deep Blue Luminescent Lead Bromide Perovskite Microdisks. Chem. Commun. 2015, 51(91), 16385–16388. https://doi.org/10.1039/C5CC06750B. [49] Filippetti, A.; Spaldin, N. A. Self-Interaction-Corrected Pseudopotential Scheme for Magnetic and Strongly-Correlated Systems. Phys. Rev. B 2003, 67(12), 125109. https://doi.org/10.1103/PhysRevB.67.125109. [50] Amgar, D.; Wierzbowska, M.; Uvarov, V.; Gutkin, V.; Etgar, L. Novel Rubidium Lead Chloride Nanocrystals: Synthesis and Characterization. Nano Futur. 2017, 1(2), 021002. https://doi.org/10.1088/2399-1984/aa8761. [51] Monzel, H.; Schramm, M.; Stöwe, K.; Beck, H. P. Zur Neuuntersuchung Des Phasendiagramms RbCl/PbCl2. Zeitschrift für Anorg. und Allg. Chemie. 2000, 626(2), 408–411. https://doi.org/10.1002/(SICI)1521-3749(200002)626:23.0.CO;2-A. [52] Ma, W.; Swisher, S. L.; Ewers, T.; Engel, J.; Ferry, V. E.; Atwater, H. A.; Alivisatos, A. P. Photovoltaic Performance of Ultrasmall PbSe Quantum Dots. ACS Nano 2011, 5(10), 8140–8147. https://doi.org/10.1021/nn202786g. [53] Pan, A.; He, B.; Fan, X.; Liu, Z.; Urban, J. J.; Alivisatos, A. P.; He, L.; Liu, Y. Insight into the Ligand-Mediated Synthesis of Colloidal CsPbBr 3 Perovskite Nanocrystals: The Role of Organic Acid, Base, and Cesium Precursors. ACS Nano 2016, 10(8), 7943–7954. https://doi.org/10.1021/acsnano.6b03863. [54] De Roo, J.; Ibáñez, M.; Geiregat, P.; Nedelcu, G.; Walravens, W.; Maes, J.; Martins, J. C.; Van Driessche, I.; Kovalenko, M. V.; Hens, Z. Highly Dynamic Ligand Binding and Light Absorption Coefficient of Cesium Lead Bromide Perovskite Nanocrystals. ACS Nano 2016, 10(2), 2071–2081. https://doi.org/10.1021/acsnano.5b06295. [55] Chen, Y.; Wu, Q.; Yin, R.; Ding, Y. Facile Fabrication and Optical Properties of Novel Pb(OH)Cl Nanotubes. J. Nanoparticle Res. 2007, 9(2), 283–287. https://doi.org/10.1007/s11051-006-9172-8.

References

57

[56] Shu, J.; Ma, R.; Shao, L.; Shui, M.; Wang, D.; Wu, K.; Long, N.; Ren, Y. Hydrothermal Fabrication of Lead Hydroxide Chloride as a Novel Anode Material for Lithium-Ion Batteries. Electrochim. Acta 2013, 102, 381–387. https://doi.org/10.1016/J. ELECTACTA.2013.04.023. [57] Woensdregt, C. F.; Hartman, P. Structural Morphology of Cotunnite, PbCl2, Laurionite, Pb(OH)Cl, and SbSI. J. Cryst. Growth 1988, 87(4), 561–566. https://doi.org/10.1016/0022-0248(88)90105-4. [58] Shen, H.; Li, J.; Shang, H.; Niu, J.; Xu, W.; Wang, H.; Guo, F.; Li, L. S. Phosphine-Free Synthesis from 1D Pb(OH)Cl Nanowires to 0D and 1D PbSe Nanocrystals. ACS Appl. Mater. Interfaces 2013, 5(20), 10331–10336. https://doi.org/10.1021/am403361z. [59] Shannon, R. D. Revised Effective Ionic Radii and Systematic Studies of Interatomic Distances in Halides and Chalcogenides. Acta Crystallogr. Sect. A 1976, 32(5), 751–767. https://doi.org/10.1107/S0567739476001551. [60] Amgar, D.; Binyamin, T.; Uvarov, V.; Etgar, L. Near Ultra-Violet to Mid-Visible Band Gap Tuning of Mixed Cation Rb x Cs 1−x PbX 3 (X = Cl or Br) Perovskite Nanoparticles. Nanoscale 2018, 10(13), 6060–6068. https://doi.org/10.1039/C7NR09607K. [61] Glazer, A. M.; IUCr. The Classification of Tilted Octahedra in Perovskites. Acta Crystallogr. Sect. B Struct. Crystallogr. Cryst. Chem. 1972, 28(11), 3384–3392. https://doi.org/10.1107/S0567740872007976. [62] Beck, H. P.; Schramm, M.; Haberkorn, R.; Dinnebier, R.; Stephens, P. W. Synthesis and Crystal Structure of Rb6Pb5Cl16. Zeitschrift für Anorg. und Allg. Chemie. 1998, 624(3), 393–398. https://doi.org/10.1002/(SICI)1521-3749(199803)624:33.0. CO;2-V. [63] Howard, C. J.; Stokes, H. T. Group-Theoretical Analysis of Octahedral Tilting in Perovskites. Acta Crystallogr. Sect. B Struct. Sci. 1998, 54(6), 782–789. https://doi.org/10.1107/S0108768198004200. [64] Beck, H. P.; Milius, W. Study on A4BX6 Compounds. II [1]. Refinement of the Structure of Rb4PbBr6 and a Note on the Existence of ?Rb4Hgl6? And ?K4CdI6? Zeitschrift für Anorg. und Allg. Chemie. 1988, 562(1), 102–104. https://doi.org/10.1002/zaac.19885620112. [65] Cola, M.; Massarotti, V.; Riccardi, R.; Sinistri, C. Binary Systems Formed by Lead Bromide with (Li, Na, K, Rb, Cs and Tl)Br: A DTA and Diffractometric Study. Zeitschrift für Naturforsch. A. 1971, 26(8), 1328–1332. https://doi.org/10.1515/zna-1971-0812. [66] Beck, H. P.; Milius, W. Study on A4BX6 Compounds. III [1]. Ns2 Cations as a Prerequisite for a Structure Type and Their Interaction in Ternary Halides with the Formula Type A4BX6 (A: In, TI; B: Cd, Pb, Ge; X: Cl, Br, I). Zeitschrift für Anorg. und Allg. Chemie. 1988, 562(1), 105–114. https://doi.org/10.1002/zaac.19885620113. [67] Meloni, S.; Palermo, G.; Astani, N. A.; Curchod, B. F. E.; Graetzel, M.; Roethlisberger, U. Valence and Conduction Bands Engineering in Halide Perovskites for Solar Cell Applica-Tions. 2014. [68] Udayabhaskararao, T.; Kazes, M.; Houben, L.; Lin, H.; Oron, D. Nucleation, Growth, and Structural Transformations of Perovskite Nanocrystals. Chem. Mater. 2017, 29(3), 1302–1308. https://doi.org/10.1021/acs.chemmater.6b04841. [69] Wu, X.; Trinh, M. T.; Zhu, X.-Y. Excitonic Many-Body Interactions in Two-Dimensional Lead Iodide Perovskite Quantum Wells. J. Phys. Chem. C. 2015, 119(26), 14714–14721. https://doi.org/10.1021/acs.jpcc.5b00148.

58

References

[70] Hassan, Y.; Song, Y.; Pensack, R. D.; Abdelrahman, A. I.; Kobayashi, Y.; Winnik, M. A.; Scholes, G. D. Structure-Tuned Lead Halide Perovskite Nanocrystals. Adv. Mater. 2016, 28(3), 566–573. https://doi.org/10.1002/adma.201503461. [71] Aharon, S.; Cohen, B. El; Etgar, L. Hybrid Lead Halide Iodide and Lead Halide Bromide in Efficient Hole Conductor Free Perovskite Solar Cell. J. Phys. Chem. C. 2014, 118(30), 17160–17165. https://doi.org/10.1021/jp5023407. [72] Aharon, S.; Etgar, L. Two Dimensional Organometal Halide Perovskite Nanorods with Tunable Optical Properties. Nano Lett. 2016, 16(5), 3230–3235. https://doi.org/10.1021/ acs.nanolett.6b00665. [73] Kangning Liang; David B. Mitzi, * and; Prikas, M. T. Synthesis and Characterization of Organic−Inorganic Perovskite Thin Films Prepared Using a Versatile Two-Step Dipping Technique. 1998. https://doi.org/10.1021/CM970568F. [74] Schmidt, L. C.; Pertegás, A.; González-Carrero, S.; Malinkiewicz, O.; Agouram, S.; Mínguez Espallargas, G.; Bolink, H. J.; Galian, R. E.; Pérez-Prieto, J. Nontemplate Synthesis of CH 3 NH 3 PbBr 3 Perovskite Nanoparticles. J. Am. Chem. Soc. 2014, 136(3), 850–853. https://doi.org/10.1021/ja4109209. [75] Dou, L.; Wong, A. B.; Yu, Y.; Lai, M.; Kornienko, N.; Eaton, S. W.; Fu, A.; Bischak, C. G.; Ma, J.; Ding, T.; et al. Atomically Thin Two-Dimensional Organic-Inorganic Hybrid Perovskites. Science 2015, 349(6255), 1518–1521. https://doi.org/10.1126/science.aac7660. [76] Kitazawa, N. Excitons in Two-Dimensional Layered Perovskite Compounds: (C6H5C2H4NH3)2Pb(Br,I)4 and (C6H5C2H4NH3)2Pb(Cl,Br)4. Mater. Sci. Eng. B 1997, 49(3), 233–238. https://doi.org/10.1016/S0921-5107(97)00132-3. [77] Ishihara, T.; Takahashi, J.; Goto, T. Optical Properties Due to Electronic Transitions in Two-Dimensional Semiconductors ( C n H 2 n + 1 NH 3) 2 PbI 4. Phys. Rev. B. 1990, 42(17), 11099–11107. https://doi.org/10.1103/PhysRevB.42.11099. [78] Ishihara, T. Optical Properties of PbI-Based Perovskite Structures. J. Lumin. 1994, 60–61, 269–274. https://doi.org/10.1016/0022-2313(94)90145-7. [79] Aharon, S.; Wierzbowska, M.; Etgar, L. The Effect of the Alkylammonium Ligand’s Length on Organic–Inorganic Perovskite Nanoparticles. ACS Energy Lett. 2018, 3(6), 1387–1393. https://doi.org/10.1021/acsenergylett.8b00554. [80] Xing, J.; Yan, F.; Zhao, Y.; Chen, S.; Yu, H.; Zhang, Q.; Zeng, R.; Demir, H. V.; Sun, X.; Huan, A.; et al. High-Efficiency Light-Emitting Diodes of Organometal Halide Perovskite Amorphous Nanoparticles. ACS Nano 2016, 10(7), 6623–6630. https://doi.org/10.1021/acsnano.6b01540. [81] Bhaumik, S.; Veldhuis, S. A.; Ng, Y. F.; Li, M.; Muduli, S. K.; Sum, T. C.; Damodaran, B.; Mhaisalkar, S.; Mathews, N. Highly Stable, Luminescent Core–Shell Type Methylammonium–Octylammonium Lead Bromide Layered Perovskite Nanoparticles. Chem. Commun. 2016, 52(44), 7118–7121. https://doi.org/10.1039/C6CC01056C. [82] Jang, D. M.; Park, K.; Kim, D. H.; Park, J.; Shojaei, F.; Kang, H. S.; Ahn, J.-P.; Lee, J. W.; Song, J. K. Reversible Halide Exchange Reaction of Organometal Trihalide Perovskite Colloidal Nanocrystals for Full-Range Band Gap Tuning. Nano Lett. 2015, 15(8), 5191–5199. https://doi.org/10.1021/acs.nanolett.5b01430. [83] Venkataraman, N. V.; Bhagyalakshmi, S.; Vasudevan, S.; Seshadri, R. Conformation and Orientation of Alkyl Chains in the Layered Organic–Inorganic Hybrids: (C n H 2n+1 NH 3) 2 PbI 4 (n = 12,16,18). Phys. Chem. Chem. Phys. 2002, 4(18), 4533–4538. https://doi.org/10.1039/B204983J.

References

59

[84] Li, C.; Wan, J.; Zheng, Y.; Dong, W. Tuning Photoconductive Properties of Organic–Inorganic Hybrid Perovskite Nanocomposite Device via Organic Layer’s Thickness. Mater. Lett. 2012, 76, 187–189. https://doi.org/10.1016/J.MATLET.2012.02.060. [85] Cao, D. H.; Stoumpos, C. C.; Farha, O. K.; Hupp, J. T.; Kanatzidis, M. G. 2D Homologous Perovskites as Light-Absorbing Materials for Solar Cell Applications. J. Am. Chem. Soc. 2015, 137(24), 7843–7850. https://doi.org/10.1021/jacs.5b03796.