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Synthesis, Self-Assembly and High-Pressure Properties of Nanoparticles and Hybrid Nanocomposites
Lingyao Meng University of New Mexico - Main Campus
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Lingyao Meng Candidate
Nanoscience and Microsystems Engineering Department
This dissertation is approved, and it is acceptable in quality and form for publication:
Approved by the Dissertation Committee:
Prof. Yang Qin, Chairperson
Prof. John Grey
Prof. Terefe Habteyes
Prof. Hongyou Fan
i Synthesis, Self-Assembly and High-Pressure Properties of Nanoparticles and Hybrid Nanocomposites
by
Lingyao Meng
M.Sc. Materials Science and Engineering, University of Pennsylvania, 2015
DISSERTATION
Submitted in Partial Fulfillment of the Requirements for the Degree of
Doctor of Philosophy Nanoscience and Microsystems Engineering
The University of New Mexico Albuquerque, New Mexico
July 2020
ii DEDICATION
-To my husband, Dejun, who has always been a constant source of love, support, help and encouragement since I was sixteen.
-To my parents, who have loved me unconditionally and believed in me more than I did myself.
iii ACKNOWLEDGEMENTS
First and foremost, I would like to thank my advisor Dr. Yang Qin, for offering me outstanding supervision and mentoring during my PhD study and research. I am grateful for his endless support, motivation, encouragement and knowledge in guiding me to become an independent researcher.
I would like to greatly acknowledge the members of my PhD committee, Dr. John
Grey, Dr. Terefe Habteyes, and Dr. Hongyou Fan for their careful revision of my dissertation, for being engaged in my research and pointing me in the right direction. I am thankful for the collaboration and useful discussions with Dr. Hongyou Fan on the research project described in Chapter 4.
I would like to thank all our collaborators as well, especially Dr. Ying-bing Jiang for the training and help with TEM and SEM, and Dr. Eric John Peterson for the XRD training. I also thank Dr. Changyong Park, Dr. Paul Chow and Dr. Yuming Xiao at the
Argonne National Laboratory for their efforts to set up and calibrate high pressure
SAXS/WAXS instruments as well as the time they took to teach me.
The past and present members of the Qin research group have provided an incalculable amount of help to me: Jianzhong, Wenhan, Keda, Zhen, Brad, Rob, Sheela,
Ellie, Yousef, Jin and Chao. They are a group of talented people, it has been a pleasure working with them. Special thanks to Zhen, for his help with instruments and always being excited to talk about scientific problems or data; to Sheela, who has always been there to talk to and hang out with when I needed a break.
I would also like to express my thank to my best friends, Lengge, Dairan, Shanshan,
Betty, Liye and Kun, for their constant support, motivation and help.
iv Finally, many thanks to my husband and my parents who have been so supportive throughout the years of my study. I could not have completed this journey without their backing and encouragement.
v Synthesis, Self-Assembly and High-Pressure Properties of Nanoparticles and Hybrid Nanocomposites
By
Lingyao Meng
M.Sc. Materials Science and Engineering, University of Pennsylvania, 2015 Ph.D. Nanoscience and Microsystems Engineering, University of New Mexico, 2020
Abstract
Nanoparticles have gained significant scientific interests owing to their unique structural dimensions, size- and shape-tunable properties, and numerous fascinating applications, from opto-electronics, sensor devices, to energy, environmental, and medical fields. Furthermore, the synergistic integration of other materials, including organic polymers, with nanoparticles provides new opportunities and strategies to obtain nanocomposites with superior properties and functionalities. While there is already significant research on the synthesis and characterizations of nanoparticles and hybrid nanocomposites, some research questions, such as how to design and control the interfacial morphology in polymer/nanoparticle hybrid nanocomposites, how to synthesize metal- organic framework (MOF) nanoparticles in well-defined and uniform sizes and shapes, and how the size and shape of nanoparticles affect their properties under high pressures, are still challenges of today. In order to tackle these challenges, this research thesis focuses on the synthesis, self-assembly and high-pressure properties of three different classes of nanoparticles or hybrid nanocomposite materials.
vi In the first part of this thesis, hybrid nanocomposites of conjugated polymers and inorganic nanoparticles are discussed, and a novel supramolecular strategy to assemble polymers and nanoparticles into stable and well-ordered core/shell composite nanofibers through the cooperation of several non-covalent interactions (hydrogen bonding, π-π interactions, etc.) is examined. By synthesizing conjugated polymers with specific functional groups (e.g. pyridine), we have successfully attached CdSe quantum dots and
Fe3O4 nanoparticles non-covalently onto the polymer nanofibers. Besides the excellent conducting property of the conjugated polymer, the resulting nanocomposites also show some added benefits, such as broader light absorption range when combined with quantum dots as well as added magnetic responsiveness when combined with iron oxide nanoparticles. Further incorporation of such composite nanofibers into organic photovoltaic devices has led to the formation of well-dispersed photoactive layer morphology. This strategy can be used as general design principles for assembling incompatible hybrid nanocomponents into well-ordered structures.
The second part of this thesis focuses on formation strategies and mechanisms for well-defined one-dimensional (1D) MOF nanostructures. Unlike inorganic and organic nanoparticles, for which synthetic procedures have well been established, generalizable design and preparation of MOF nanoparticles are still under early developing stages. To address this challenge, we have developed a new method to rapidly and reproducibly synthesize continuous 1-D MOF nano/micro-structures through interfacial synthesis templated by nanoporous polymer membranes. In this study, zeolitic imidazole frameworks (ZIFs) and polycarbonate track-etched (PCTE) membranes were used as model materials, and by varying the experimental conditions (pore sizes, reaction time,
vii metal ion source, etc.), different 1D ZIF-8 or ZIF-67 nano/micro-structures with the pore dimensions corresponding to the PCTE templates were obtained, which were further fully characterized by a combination of electron microscopy and X-ray techniques. This work represents the first example of membrane templated synthesis of MOF nano/micro- structures. Our findings provide a generalized method for controlling size, morphology, and lattice orientation of MOF nanomaterials.
The last part of this thesis discusses how the size and shape of nanoparticles influence their pressure-dependent properties using CdS nanoparticles as a model material.
CdS nanoparticles are synthesized in three different sizes and shapes, and are subjected to controlled high pressures up to 15 GPa in a diamond anvil cell. Characterizations with in- situ small and wide-angle X-ray scattering measurements under high pressure suggest that both the reversibility of phase transition and phase transition pressure are closely related to the particle size and shape. Further characterizations with transmission electron microscopy show that external pressure can decrease the nanoparticle separation distance and induce sintering and coalescence of nanoparticles into new nanostructures. Our results provide new insights into the fundamental properties of nanoparticles under high pressure that will inform designs of new nanomaterial structures for emerging applications.
viii Table of Contents
ACKNOWLEDGEMENTS ...... iv
Abstract ...... vi
List of Figures ...... xii
List of Tables ...... xvii
Chapter 1. Introduction ...... 1
1.1 Overview ...... 1
1.2 Nanoparticles ...... 2
1.2.1 Size- and Shape- Dependent Properties of Nanoparticles ...... 3
1.2.2 Synthesis of Nanoparticles ...... 5
1.2.3 Self-Assembly of Nanoparticles ...... 6
1.3 Nanocomposites of Polymer and Inorganic Nanoparticles ...... 7
1.3.1 Polymer/Nanoparticle Nanocomposites Synthesis ...... 8
1.3.2 Polymer/Nanoparticle Nanocomposites for Photovoltaic Applications ...... 10
1.4 Metal-Organic Framework Nanoparticles ...... 12
1.4.1 Metal-Organic Frameworks ...... 13
1.4.2 Metal-Organic Framework Nanoparticles ...... 13
1.5 Nanoparticles under Pressure ...... 15
1.5.1 High Pressure Characterization Techniques ...... 15
1.5.2 Structures of Nanoparticles under Pressure ...... 16
1.5.3 Properties and Applications of Nanoparticles under Pressure ...... 17
1.5.4 CdS Nanoparticles under Pressure ...... 19
1.6 Motivations of My Projects ...... 21
ix 1.7 References ...... 22
Chapter 2. Bottom-Up Approach for Precisely Nanostructuring Hybrid
Organic/Inorganic Multicomponent Composites for Organic Photovoltaics ...... 44
2.1 Introduction ...... 44
2.2 Synthesis and Characterization ...... 47
2.2.1 Synthetic Procedures ...... 47
2.2.2 Characterizations ...... 51
2.2.3 Solar Cell Fabrication and Measurement ...... 52
2.3 Results and Discussions ...... 53
2.3.1 Synthesis of Polymer Nanofibers and Inorganic Nanoparticles ...... 53
2.3.2 Self-Assembly of Hybrid Conjugated Polymer/Quantum Dot Composite
Nanofibers ...... 58
2.3.3 Organic Solar Cells Fabricated from Hybrid Conjugated Polymer/Quantum Dot
Composite Nanofibers ...... 60
2.3.4 Self-Assembly and Magnetic Responses of Hybrid Conjugated
Polymer/Magnetic Nanoparticle Composite Nanofibers ...... 61
2.3.5 Organic Solar Cells Fabricated from Hybrid Conjugated Polymer/Magnetic
Nanoparticle Composite Nanofibers and the Active Layer Morphology ...... 68
2.4 Conclusions ...... 72
2.5 References ...... 73
Chapter 3. Metal-Organic Framework (MOF) One-Dimensional Nanostructures .. 84
3.1 Introduction ...... 84
3.2 Synthesis and Characterization ...... 86
x 3.2.1 Synthetic Procedures ...... 86
3.2.2 Characterizations ...... 87
3.3 Results and Discussions ...... 88
3.4 Conclusions ...... 100
3.5 References ...... 100
Chapter 4. Size and Shape Dependence of Pressure Induced Phase Transition in CdS
Semiconductor Nanocrystals ...... 108
4.1 Introduction ...... 108
4.2 Synthesis and Characterization ...... 111
4.2.1 Synthetic Procedures ...... 111
4.2.2 Characterizations ...... 114
4.3 Results and Discussions ...... 115
4.3.1 Size Dependence of Pressure-Induced Phase Transition in CdS Semiconductor
Nanocrystals ...... 115
4.3.2 Shape Dependence of Pressure-Induced Phase Transition in CdS Semiconductor
Nanocrystals ...... 121
4.4 Conclusions ...... 129
4.5 References ...... 130
xi List of Figures Chapter 1
Figure 1. 1 Size- and shape- dependent properties of nanoparticles ...... 3
Figure 1. 2 Hot injection synthesis of nanoparticles ...... 5
Figure 1. 3 Self-assembly of nanoparticles ...... 6
Figure 1. 4 Schemes showing the main chemical routes for the synthesis of polymer/nanoparticle nanocomposites ...... 9
Figure 1. 5 The bulk heterojunction structure ...... 10
Figure 1. 6 Synthesis of MOF nanoparticles ...... 13
Figure 1. 7 Photo and schematic illustration of a diamond anvil cell ...... 15
Figure 1. 8 Scheme illustrating the outline of this thesis...... 22
Chapter 2
Figure 2. 1 Synthetic scheme of block copolymers and CdSe QDs capping ligands...... 53
Figure 2. 2 UV-vis absorption spectra of nanofiber solutions of (A) P3HT, (B) BCP2, and
(C) BCP3...... 55
Figure 2. 3 Transmission electron microscopy (TEM) images of (A) P3HT NFs; (B) BCP2
NFs; (C) BCP3 NFs; (D) IONP-OA; (E) IONP-L-OA; and (F) IONP-CA. Inserts: histograms of corresponding NF widths and nanoparticle diameters sampled from 100 individual objects...... 56
Figure 2. 4 TEM images of (A) CdSe QDs with TOPO ligands (B) CdSe QDs with PDTC ligands...... 57
Figure 2. 5 Selected area electron diffraction (SAED) pattern (A) and powder X-ray diffraction spectrum (B) of IONP-OA...... 58
xii Figure 2. 6 TEM images of (A) P3HT NFs with CdSe QDs having TOPO ligands (1/1, w/w); (B) BCP3 NFs with CdSe QDs having TOPO ligands (1/1, w/w); (C) P3HT NFs with
CdSe QDs having PDTC ligands (1/1, w/w); (D) BCP3 NFs with CdSe QDs having PDTC ligands (1/1, w/w)...... 60
Figure 2. 7 TEM images of nanostructures from mixtures of P3HT NFs and (A) IONP-
OA, (B) IONP-L-OA, and (C) IONP-CA; BCP2 NFs and (D) IONP-OA, (E) IONP-L-OA, and (F) IONP-CA; and BCP3 NFs and (G) IONP-OA, (H) IONP-L-OA, and (I) IONP-CA.
The mixture solutions used for TEM analyses contain polymer NFs and IONPs at a ca. 2/1 weight ratio and polymer concentrations at ca. 0.1 mg mL−1. Scale bars in all: 200 nm. . 62
Figure 2. 8 Infrared (IR) spectra on powders of IONP-OA (black), BCP3 (Blue), and precipitate of BCP3/IONP-OA mixture (red)...... 65
Figure 2. 9 Photographs of solutions of composite NFs next to a permanent magnetic cube at the start time and the times when solutions became mostly clear. Durations for such processes to take place are shown above arrows (s: second; m: minute; h: hour)...... 65
Figure 2. 10 Photographs of well-dissolved solutions of P3HT (left), BCP2 (middle), and
BCP3 (right) mixed with IONP-OA in chlorobenzene (2/1 wt./wt., 10 mg/mL polymer concentration) placed next to a permanent magnet cube. In each photo, the mixture solutions sit on the left and on the right are the pure IONP-OA solutions in chlorobenzene at identical concentrations...... 67
Figure 2. 11 TEM images of device active layers employing (A) P3HT NF/PCBM; (B)
P3HT NF/PCBM/IONP-OA; (C) BCP3 NF/PCBM; and (D) BCP3 NF/PCBM/IONP-OA.
Scale bars in all: 200 nm...... 70
xiii Figure 2. 12 X-ray diffraction (XRD) profiles of thin films (100 nm in thickness, thermal annealed at 150 °C for 10 min) of BCP3 NF (black), BCP3 NF/PCBM (red), BCP3
NF/IONP (blue) and BCP3 NF/IONP/PCBM (green)...... 72
Chapter 3
Figure 3. 1 Image of the reaction vial for the synthesis of ZIF-8 1D nanostructures, the dotted circle represents the PCTE membrane...... 89
Figure 3. 2 XRD patterns of as-synthesized (A) ZIF-8 and (B) ZIF-67 membranes with different pore sizes...... 90
Figure 3. 3 SEM images of both the hydrophobic and hydrophilic surfaces of ZIF-8 and
ZIF-67 containing PCTE membranes with different pore size. The scales of each image are shown on the left...... 93
Figure 3. 4 (a)-(j) TEM images of ZIF-8 and ZIF-67 nanowires and nanorods formed within the pores of the PCTE membranes. Insets in (a)-(h) are TEM images with lower magnification, insets in (i) and (j) are optical microscopic images; (k)-(n) SEM images of
ZIF-8 and ZIF-67 nanocylinders and nanodisks formed within the pores of the PCTE membranes. Insets in (k)-(n) are optical microscopic images...... 95
Figure 3. 5 SEM images of isolated nanowires, nanorods, nanocylinders and nanodisks.
The scales of each image are shown on the left...... 97
Figure 3. 6 TEM images and the corresponding SAED patterns of (A) (C) ZIF-67 100 nm nanorods and (B) (D) ZIF-67 30 nm nanorods. Samples were calibrated with silicon standard. The red circle shows where the SAED pattern was obtained...... 99
Chapter 4
xiv Figure 4. 1 TEM images of (a) 7.5 nm, (b) 10.6 nm, (c) 39.7 nm CdS nanoparticles and (d) corresponding room pressure powder XRD spectrum...... 115
Figure 4. 2 Representative synchrotron WAXS data during compression and decompression of (a) 7.5 nm, (b) 10.6 nm and (c) 39.7 nm CdS nanoparticles. r represents the releasing pressure, the black curve represents the WZ phase, the blue curve represents the RS phase, and the red curve represents a mixture of WZ and RS. Impurity peaks from gasket Re, neon gas and ruby are marked with asterisk...... 117
Figure 4. 3 TEM images of (a) 7.5 nm, (b) 10.6 nm, and (c) 39.7 nm CdS nanoparticle samples after compression and decompression process...... 118
Figure 4. 4 Dependence of unit cell volume on the applied external pressure for (a) 7.5, (b)
10.6, and (c) 39.7 nm CdS nanoparticles. Black dots represent the compression process and red dots represent the decompression process...... 120
Figure 4. 5 Transmission electron microscopy (TEM) images of (a) spherical CdS nanoparticles, (b) short CdS nanorods, and (c) long CdSe/CdS core/shell nanorods. .... 121
Figure 4. 6 Wide-angle X-ray scattering (WAXS) patterns under various applied pressure:
(a) CdS nanospheres, (b) short CdS nanorods and (c) long CdSe/CdS core/shell nanorods; during compression and decompression. Pressures labeled with letter r are during decompression processes. The black, blue, and red curves represent the WZ, RS, and
WZ/RS mixture crystal structures, respectively. The red asterisks mark diffraction peaks from rhenium gaskets used in the anvil cells...... 124
Figure 4. 7 TEM images of (a) CdS nanospheres, (b) short CdS nanorods, (c) long
CdSe/CdS core/shell nanorods after high pressure studies; and high-resolution TEM (HR-
xv TEM) images of (d) CdS nanospheres, (e) short CdS nanorods, and (f) long CdSe/CdS core/shell nanorods after high pressure studies...... 125
Figure 4. 8 TEM images of spherical CdS nanoparticles after compression...... 126
Figure 4. 9 TEM images of (a) short CdS nanorods and (b) long CdS nanorods after compression...... 127
Figure 4. 10 Pressure dependence of the unit cell volume for (a) CdS nanospheres, (b) short CdS nanorods, and (c) long CdSe/CdS core/shell nanorods. The black and red dots represent the compression and decompression process, respectively...... 129
xvi List of Tables
Chapter 2
Table 2. 1 Summary of Solar Cell Device Performance Data ...... 60
Table 2. 2 Average numbers (Navg) and maximum numbers (Nmax) of IONPs closely associated with one polymer NF, from sampling ca. 50 individual NFs in TEM images . 61
Table 2. 3 Organic solar cell performance parameters using P3HT and BCP3 NFs in combination with PCBM and varied amount of IONP-OA ...... 68
Chapter 3
Table 3. 1 Reactions conditions for different sizes MOF nanostructures...... 87
Table 3. 2 Pore size and thickness of the PCTE membrane...... 88
Table 3. 3 CPO indices for different sizes MOF nanostructures...... 91
Table 3. 4 Diameters and lengths of MOF nanostructures ...... 98
Chapter 4
Table 4. 1 Unit cell volumes and bulk moduli of CdS nanoparticles...... 120
Table 4. 2 Size and surface-to-volume ratio of CdS nanoparticles...... 121
Table 4. 3 Calculated unit cell volumes and bulk moduli of the three CdS samples...... 127
xvii Chapter 1. Introduction 1.1 Overview
Nanotechnology is an emerging field that focuses on the design, synthesis and fabrication of materials or structures at the nanometer scales and can be applied through all other scientific fields such as chemistry, material science, physics, biology and engineering.1-2 Nanomaterials are the foundation of nanotechnology with a broad range of applications like electronics,3 optics,4 sensors,5 information storage,6 energy conversion7 and medicine.8 A number of techniques, from complex synthesis, to X-ray, spectroscopy and scattering techniques have been used to investigate nanomaterials’ characteristics.9-10
So far, the properties of nanomaterials discovered have been remarkable, for example, semiconducting quantum dots (QDs) and metal nanoparticles show size- and shape- tunable properties,11 carbon nanotubes possess excellent mechanical, thermal and conductive properties,12 electrospun nanofibers show superhydrophobic surface properties.13 With coherent advances in the synthesis, experimental and theoretical understanding of nanomaterials, the potential of the nanoworld seems truly unlimited.
Although the properties of an individual nanomaterial are interesting, they are still limited by their constituent elements and compositions. The creation of a composite material makes the combination of the properties of different nanomaterials possible.
However, the performances of the nanocomposites depend not only on the intrinsic properties of isolated nanomaterials, but also on their spatial arrangements or ordering.
Consequently, the understanding and controlling of materials interfaces are of great importance in the design and fabrication of novel nanocomposites. On the other hand, many current strategies for nanomaterial synthesis and characterizations are developed at ambient conditions. The high-pressure behaviors of nanomaterials have been relatively less
1 explored, and many of the nanoparticle high-pressure properties are still not well understood. In this thesis, research focused on solving those challenges are presented. The first two parts of this thesis discuss two different approaches to form stable and controlled organic/inorganic hybrid nanostructures with precisely controlled hybrid domain shapes and sizes. The third part focuses on investigating nanoparticles phase properties under high pressure.
This introductory chapter aims to provide a background and outline for the research.
As the separate projects covered in this thesis differ substantially in some areas, this chapter provides only the briefest of introductions, and more detailed discussions will be provided in the specific chapter they relate to.
1.2 Nanoparticles
The word “nano” comes from ancient Greek, meaning “dwarf”. A nanometer (nm) is one billionth of a meter, or 10-9 m. Materials with at least one of the three external dimensions in the nanoscale regime (1-100 nm) can be referred to as nanoscaled.14
Nanomaterials can be classified based on the number of dimensions they display in the nanoscale.15 Nanocoils, nanopillars, nanocones, nanoballs and nanoflowers are known as three-dimensional (3D) nanomaterials. The two-dimensional 2D nanostructures generally have their length and width larger than their thickness. Materials such as graphene and graphene oxide are the most well-known 2D nanomaterials. The one-dimensional (1D) nanomaterials have one dimension much greater than the others. Common structures include nanofibers, nanotubes, nanorods and nanowires. The zero-dimensional (0D) nanomaterials refer to structures having all of its dimensions in the nanoscale and are simply called nanoparticles. The nanoparticles can be of different types comprising carbon-
2 based nanoparticles, ceramic nanoparticles, metal or metal oxide nanoparticles, quantum dots, polymeric nanoparticles and lipid-based nanoparticles.16 Nowadays, nanoparticles have attracted a great deal of research interests because they can bridge the gap between the physics/chemistry of atomic structures and the bulk materials.
1.2.1 Size- and Shape- Dependent Properties of Nanoparticles
Figure 1. 1 Size- and shape- dependent properties of nanoparticles: (A) Bright fluorescence from CdSe/ZnS core/shell quantum dots with CdSe core diameters ranging from 6.9 nm to
1.8 nm (adapted from Ref. 17); (B) Bandgap energy increases as the quantum dot size decreases (adapted from Ref. 18); (C) Comparison of hysteresis loops of superparamagnetic nanoparticles and ferromagnetic nanoparticles (adapted from Ref. 19);
(D) Shift of the bandgap of CdSe nanoparticles confined in 3D (QDs), 2D (quantum rods), or 1D (quantum disks) (adapted from Ref. 20); (E) Magnetic hysteresis curves of (a) solid
Fe3O4 nanospheres, (b) solid Fe3O4 nanoellipsoids, (c) hollow Fe3O4 nanoellipsoids, inset shows the low field region magnification (adapted from Ref. 21).
3 At small sizes, the number of surface atoms becomes a significant fraction of the total number of atoms in the crystallite and the surface energy plays an important role in determining its thermal stability. Numerous material properties can thus be affected by decreasing the grain size, and many novel size-dependent properties of nanoparticles have already been discovered. For example, the melting temperature of a nanoparticle decreases with decreasing the particle size due to the increase of the surface energy.22-23 Mechanical properties, such as elastic modulus,24-26 also vary with particle size. In addition, QDs, which are semiconductor nanoparticles, are well-known for their size-dependent optical properties and bandgaps.27-28 As the particle size decreases, their bandgap increases and the absorption wavelength shift to the blue region (Figures 1.1 (A)-(B)), which can be explained by the concept of the quantum confinement effect. In the case of magnetic nanoparticles, as the size of ferromagnetic or ferrimagnetic nanoparticles decrease, they can act as a single magnetic domain and their magnetic moments are affected by thermal energy.29 Such particle is said to be superparamagnetic (Figure 1.1 (C)). Besides the particle size, nanoparticle properties are also dependent on their shape.30 For example,
Figures 1.1 (D)-(E) show that both the bandgap of QDs, and the saturation magnetization values in magnetic nanoparticles are shape-dependent. These unique properties have made nanoparticles popular for diverse applications such as photovoltaic devices,31 field effect transistors,32 biosensors33 and bio-imaging.34
4 1.2.2 Synthesis of Nanoparticles
Figure 1. 2 Hot injection synthesis of nanoparticles (adapted from Ref. 36-37): (A)
Schematic of the nucleation and growth stage for the preparation of monodispersed nanoparticles, (B) Synthetic apparatus employed to prepare monodispersed nanoparticle.
The preparation of nanoscale materials can be achieved through either the “bottom- up” or the “top-down” approach.35 For top-down approaches, nanoobjects are constructed from larger entities. One of the most used top-down approaches is lithography. The major downside of this approach is the surface imperfection of the resulting nanomaterials, it is also well-known that lithography can cause significant damage to crystal structures. On the other hand, for bottom-up approaches, small building blocks are assembled into larger structures, which makes the formation of homogeneous nanostructures with less defects possible. The hot-injection method is one of the most widely used bottom-up technique for growing monodispersed nanocrystals (Figure 1.2).36-37 It involves the rapid injection of organometallic precursor reagents into a hot solution that contains surfactant molecules/ligands which prevent nanoparticles from agglomeration. Ligands typically employed are alkylphosphine, alkylphosphine oxides, long-chain carboxylic acids and long-chain amines.27 Once nucleated, diffusion-controlled growth takes place. As the
5 growth proceeds, Ostwald Ripening occurs, in which the large particles continue to grow at the expenses of small ones.38-39 This method offers a high-level of control over the size and shape of individual particles, and the size of the resulting particles can be easily controlled by varying the temperature, reaction time, and concentration of surfactants or precursors. Additionally, this method is general and has been proven to be effective at synthesizing different kinds of nanoparticles.
1.2.3 Self-Assembly of Nanoparticles
Figure 1. 3 Self-assembly of nanoparticles (adapted from Ref. 40): (A) Self-assembly process involves control over several length scales; (B) Commonly used experimental approaches.
The technology of assembling nanoparticles into nano/micro periodic structures have attracted much attention as it can be used as an effective tool to generate functional materials and novel devices.41 Among various techniques, self-assembly has been considered as one of the most effective and simple process. Self-assembly refers to the process that nanocrystals spontaneously form an arrangement due to direct specific interactions (e.g. van der Waals interactions, electrostatic interactions), or indirectly through their environment (e.g. electromagnetic field) (Figure 1.3).42 So far, many different
6 approaches have been adopted to build highly ordered nanoparticles arrays (Figure 1.3
(B)), the simplest one is to drop cast a colloidal nanoparticle solution onto a solid substrate and allow it to dry slowly. Through such approaches, nanoparticles can self-assemble into various crystal structures including face centered cubic (FCC), body centered cubic (BCC) and hexagonal close packed (HCP) structures, just like how atoms or ions bind together to from crystalline structures.36, 43 Even more interesting packing structures can be created through the use of two types of nanoparticles.44 The resulting binary nanoparticle superlattices exhibit not only structural diversity, but also combined properties of individual components with new functionalities. To date, self-assembled nanoparticle arrays have been applied to various fields, such as photonic crystals, dye-sensitized solar cells, biosensors, etc.40
1.3 Nanocomposites of Polymer and Inorganic Nanoparticles
Besides the investigation of synthesis and self-assembly of nanoparticles, in the last few decades, many researches have been attracted towards nanocomposite materials to achieve substantial enhancements of nanoparticle properties. In addition, the nanocomposite formation can be reversible and the nanocomposite itself can also act as a building block to form even more complicated structures. Specifically, the fabrication of organic/inorganic hybrid nanocomposites comprised of polymers and inorganic nanoparticles has been widely reported in the scientific literatures as a method to combine the advantages of both classes of materials into a single entity, and to obtain new properties that can act collectively and can be enhanced relative to that of individual building blocks.45-47
7 1.3.1 Polymer/Nanoparticle Nanocomposites Synthesis
As shown in Figure 1.4, many methods have been developed to prepare the composites of organic polymers and inorganic nanoparticles.45 The main synthesis routes can be divided into four paths: (A) conventional sol-gel chemistry;48-49 (B) assembling or dispersion of well-defined nanobuilding blocks that are capped with polymerizable ligands or functional dendrimers;50 (C) self-assembly procedures by using polymer assemblies as skeleton in solution, through either “in-situ” growth, or direct attachment of nanoparticles;
50-51 and (D) integrative synthesis, such as micro-modeling methods to achieve hierarchical hybrid architectures.52 Among various methods, self-assembly procedures are the most straightforward strategies. The in-situ growth approach is the most commonly employed route for the formation of polymer/nanoparticle nanocomposites as the reduction of the metal ions into nanoparticles will not affect polymer structures. However, this method usually does not produce nanoparticles with uniform sizes and shapes.
On the other hand, pre-synthesized inorganic nanoparticles can be directly attached onto polymers through complementary interactions between them. The spatial arrangements of nanoparticles are dependent on the complementary sites that are generally designed at the outer shell of the assemblies. This self-assemble method is usually applicable to host relatively small nanoparticles, such as quantum dots.53 Both covalent interactions,54-60 and a number of non-covalent interactions, such as, hydrogen bonding,61 hydrophobic interactions,62-65 and coordination interactions66 have all been employed as critical driving forces in this self-assembly process and have led to a number of interesting nanocomposite structures. The properties of resulting nanocomposites depend on a number of nanoparticle or polymer features, such as the size, aspect ratio of the nanoparticle, and
8 the molecular weight, polydispersity of the polymer. In addition, the nanoparticle/polymer interface as well as the dispersion state of nanoparticles are essential for obtaining properties enhancement to their full potential. However, due to nonadditivity of nanoparticle interactions,67 heterogeneous composites with large aggregates of particles are normally conducted. This limitation has become one of the major barriers to achieve property enhancements in polymer/nanoparticle nanocomposites. Therefore, precisely controlling the interactions between nanoparticles and polymers to improve nanoparticle dispersion quality and structure stability has long been identified as one of the key challenges in this field. Several strategies, including either chemical approaches such as surface modifications,68-69 or physical approaches such as the application of ultrasonic vibrations,70 have all been applied to achieve uniformly dispersed nanoparticles in the polymer matrix.
Figure 1. 4 Schemes showing the main chemical routes for the synthesis of polymer/nanoparticle nanocomposites (adapted from Ref. 45).
9 1.3.2 Polymer/Nanoparticle Nanocomposites for Photovoltaic Applications
The ability to incorporate nanoparticles into polymer matrices is critically important to a number of applications including electronics,71 optics,72 batteries,73 membrane separation,74 automotive parts,75 biomedical applications,76 etc.45, 77 In recent years, as the global energy demand grows, polymer/nanoparticle nanocomposites for photovoltaic (PV) applications have been widely investigated.78-80
PV is the technology that harvests energy directly from the sunlight and has been widely recognized as an essential componence of future global energy production. Over the years, PV devices can be made from different inorganic semiconductor materials, like
Si, CdTe, GaAs, and copper indium gallium selenide (CIGS), and the best crystalline silicon solar cells can reach more than 20% power conversion efficiencies (PCEs).81
However, high purity crystalline silicon wafers are fragile and very expensive.
Figure 1. 5 The bulk heterojunction structure (adapted from Ref. 82).
On the other hand, polymer solar cells (PSCs) comprising conjugated polymers
(CPs) donor and fullerene derivative acceptors can be made lightweight, flexible and possible for mass production through roll-to-roll processing.79, 83-86 The most successful strategy to date to split excitons in PSCs is by forming a bulk heterojunction (BHJ) structure (Figure 1.5).87-88 In this case, acceptor and donor materials are mixed together such that both components are intertwined with each other while remaining interconnected to either side of the device. The most commonly studied PSCs are made from BHJ blends
10 of poly(3-hexylthiophene) (P3HT) and [6,6]-phenyl-C61-butyric acid methyl ester
(PCBM).89
However, PSCs suffer from insufficient light absorption of the organic materials, as well as structural instability. In an attempt to develop a PV technology where these issues are reduced, BHJ based on CPs and semiconductor nanoparticles were devised. In principle, incorporating semiconductor nanoparticles could achieve more complete coverage of the solar spectrum and facilitate charge dissociation and transport owing to their size tunable bandgap, large surface area, relatively high dielectric constant and excellent intrinsic carrier mobilities. So far, a variety of semiconductor nanoparticles have been investigated for applications in polymer/nanoparticle hybrid solar cells, for example,
90 91 92 93 94 CdSe, PbS, PbSe, CuInS2, and TiO2. Though, these hybrid photovoltaic devices generally show efficiencies less than 5%.
The main difficult of these hybrid PV devices is that the simple blending of nanoparticles with CPs often results in a microscale phase separation, which makes the formation of well-defined hybrid interface hard to achieve. Additionally, the synthesized nanoparticles are capped with long chain organic ligands at the surface, which works as an electrical barrier. As a result, dissociation of excitons into free charge carriers at the polymer/nanoparticle interfaces will be difficult. Therefore, methodologies that could control the polymer/nanoparticle interface morphology and improve the optoelectronic interactions between them are highly desirable. Strategies described in the previous section regarding the formation of ordered polymer/nanoparticle nanocomposites have also been applied to improve the BHJ morphology of polymer/nanoparticle solar cells. However, a generalized approach to disperse nanoparticles into polymer matrix individually and
11 uniformly to form device favorable nanostructures is still missing. To date, most of them rely on strong covalent bonding, but the particle loading ratios are generally low due to the limited solubility and strong aggregation tendency of nanoparticles.59, 95-96 On the other hand, the use of non-covalent interactions (e.g. hydrogen bonding, π-π stacking), which allows for the adjustment of nanoparticle loading more easily while maintaining strong segregation, have been rarely reported. Our group has recently developed a modified solution strategy to assemble two incompatible components into well-ordered core/shell composite nanofibers (NFs) through non-covalent interactions which provides nanocomposites more homogenous and enhances the overall complex solubility. Our previous work demonstrated that through the crystallization-driven self-assembly of polythiophene based conjugated block copolymer (BCPs), fullerene derivatives could be attached onto the self-assembled BCP NFs and improved solar cell BHJ morphology and device thermal stability were achieved.97-100 In this thesis, we present the self-assembly of
BCP NFs with other types of nanoparticles, such as quantum dots and magnetic nanoparticles by implementing this composite NF strategy for PV applications, details will be discussed in Chapter 2.
1.4 Metal-Organic Framework Nanoparticles
Metal-organic frameworks (MOFs), also known as porous coordination polymers
(PCPs), have attracted increasing research interests due to their high surface areas, high porosity, well-defined 3D architectures and controllable functionalities.101-103 Even though the vast majority of research efforts have focused on MOF materials in bulk, the formation of MOF nanoparticles has also been studied recently and some unique properties have been revealed.
12 1.4.1 Metal-Organic Frameworks
MOFs are crystalline hybrid porous solids composed of a 3D network of tunable organic linkers and inorganic metal nodes (or metal cluster nodes). MOFs possess several attractive features such as high surface area, high porosity, uniform and tunable pore size, well-defined pore structures and great structural diversity. So far, a wide variety of MOFs have been reported and used in catalysis,104 gas storage,105 adsorption and separation,106 sensors,107 electronics,108 and medicine.109 Zeolite imidazolate frameworks (ZIFs) are a unique type of MOF that are composed of tetrahedrally-coordinated transition metal ions
(e.g. Fe, Co, Cu, or Zn) and imidazolate-related linkers.110 Among various ZIFs, ZIF-8
(Zn(MIM)2, MIM=2-methylimidazolate) is undoubtedly the most widely studied one as they are easily prepared, possess unexpected thermal and chemical stability, and exhibit high porosity.111 These properties make ZIF-8 ideal candidates for applications such as a porous component in mixed matrix membranes for gas separation,112-113 or as a platform to support metal nanoparticles or biomolecules to form composite materials.114-115 Besides
ZIF-8, other types of ZIF structures have also been researched for several applications, such as gas adsorption/storage, separation, electrochemistry, catalysis and so on.111, 116
1.4.2 Metal-Organic Framework Nanoparticles
Figure 1. 6 Synthesis of MOF nanoparticles (adapted from Ref. 117).
13 Unlike inorganic nanoparticles, for which synthesis methodology is very well established and a large variety of particles can be routinely synthesized, the design of generalizable MOF nanoparticle synthesis methodology is still under its early developing stage. Only recently, intensive efforts have been devoted to the preparation of MOF nanoparticles that are uniform in both size and shape (Figure 1.6).117 So far, synthesis of uniform MOF nanoparticles can be mainly divided into four approaches: (1) rapid nucleation through fast precipitation or accelerated heating;118 (2) nanoreactor confinement strategies, which isolate the nucleation site in a physical confined space;119 (3) coordination modulation via chemically controlling ligand-metal interactions;120 (4) templating methods through controlled deposition of precursor materials onto various templates, followed by selective removal of templates.121 Among those approaches, templating methods are considered as the most straightforward route.122 Accordingly, a variety of nanostructures, from 0D to 3D have all been prepared and some interesting properties that have not been observed from the bulk system such as accelerated adsorption/desorption kinetics and improved biostability, have been discovered.123-124 However, despite the tremendous progress that has been made in MOF nanoparticle synthesis, there still remains a key challenge, which is to find a reliable, general and reproducible preparation method for forming uniform MOF nanoparticles. One of the reasons behind this difficulty is that the thermodynamics of MOF formation depends on different metal-ligand combinations, thus it’s difficult to find a general synthetic method. The other reason is that the coordination bonds that drives MOF nanoparticle formation is typically weak, which makes the nanoparticle growth process slower than the precursor diffusion rate in solution, leading to a broad particle size distribution.125-126 As MOF nanoparticles can be a promising new class
14 of functional materials, further investigations on strategies to improve the MOF nanoparticle synthesis and functionalization are highly desirable. A novel strategy to synthesize 1D ZIF-8 and ZIF-67 nanostructures through interfacial synthesis templated by nanoporous polymer membranes is presented in Chapter 3.
1.5 Nanoparticles under Pressure
High pressure properties of materials have been extensively studied since the invention of the diamond anvil cell (DAC) in 1959.127-128 Because the pressure dependence of the Gibbs free energy for a homogenous material at constant temperature and composition is solely given be the volume (G = E + PV – TS), placing a material under a high pressure environment is the easiest way to examine the effects of the volume on the thermodynamics of the system. The structural phase transitions as well as the mechanical, optical, and electric properties of various materials have been investigated. Compared to bulk materials, nanoparticles are of our particular interests as their high-pressure behaviors are size and shape dependent.
1.5.1 High Pressure Characterization Techniques
Figure 1. 7 Photo and schematic illustration of a diamond anvil cell (adapted from Ref.
129).
15 The DAC (Figure 1.7) is the most commonly used device to achieve very high pressure (up to 150 GPa) in a controlled manner. A small piece of sample and a standard reference material (typically a piece of ruby) are placed in the sample chamber with a rhenium gasket, which has a hole drilled in the middle, around the side. Pressure transmitting medium, such as silicon oil or neon gas is generally added to achieve a uniform compression of the sample. Pressure is applied by pushing the two diamonds towards the sample, and it is monitored by a standard pressure-dependent ruby fluorescence technique.
As diamonds are transparency from IR to X-ray, various optical or structural techniques, including absorption, fluorescence, Raman spectroscopy and X-ray scattering experiments can be used to monitor the high-pressure behaviors of sample materials in-situ.
1.5.2 Structures of Nanoparticles under Pressure
In recent years, studies of nanomaterials under high pressure have received increasing attention because high pressure has proven to be a powerful tool to investigate the structural transformation of nanomaterials at both the atomic scale and the mesoscale.130 The structural phase transition process is generally measured by in-situ synchrotron wide-angle and small-angle X-ray scattering (WAXS or SAXS) techniques.
At atomic scale, high pressure can induce structural phase transition of materials from the ambient lattice structure to high-pressure phases. For example, CdSe bulk materials transform from the wurtzite (WZ) to the rocksalt (RS) phase at around 2.0-3.0
GPa;131-133 the phase transition pressure of PbS bulk materials from WZ to RS is at ~2.6
GPa.134-135 Nanoparticles exhibit similar phase transitions as their bulk materials. However, because the nanoparticle surface is covered by ligand molecules, which provides additional pressure shielding and minimizes the deviatoric stress, such pressure induced phase
16 transition in nanomaterials generally occurs at higher pressure. In addition, compared to the bulk material, nanoparticles possess much higher surface area to volume ratio. The surface atoms make up a significant number of the total atoms in the nanoparticle crystal, which added an additional surface term to the free energy equation. The inclusion of the surface energy also makes the phase transition of nanoparticles size- and shape- dependent.
For instance, the CdSe nanoparticles with a radius of 1 nm transformed from WZ to RS at
4.9 GPa, while the phase transition pressure of CdSe nanoparticles with a radius of 2 nm is 3.6 GPa;136-137 For 7 nm PbS nanoparticles, the phase transition pressure was reported to be 12.5 GPa.138 Additionally, it has been theoretically predicted that nanorods should have the same dependence of phase transition pressure based on the length of the nanorods.139
However, how the shape of the particle influences the high-pressure phase transition has been rarely experimentally scrutinized so far.
At mesoscale, pressure can push neighboring particles together along the pressure applied direction, nanoparticles can then contact each other and consolidate into new structures by nanoparticle coalescence. Previous research has revealed that both metal and semiconductor spherical nanoparticles can sinter into nanowires under high pressure.140-148
Other nanostructures formed by nanoparticle sintering, such as nanosheets,149-150 have also been reported. Such morphology control at the mesoscale under high pressure opens up new doors toward the synthesis of nanostructures that are difficult to obtain through solution synthesis methods.
1.5.3 Properties and Applications of Nanoparticles under Pressure
Achieving a better understanding of the high-pressure behaviors of nanoparticles can reveal valuable insights into a variety of material properties such as mechanical, optical
17 and electronic properties; this knowledge will be useful in developing future multifunctional devices that can be used in high-pressure environments.
Mechanical properties such as bulk modulus of nanomaterials can be calculated from high pressure WAXS spectra, which are generally higher than the value of the bulk material and are also size-dependent. To date, the size-dependent bulk modulus of nanoparticles has already been reported for metal,151 metal oxides,152 and semiconductors.153 However, the reported results are quite controversial: with decreasing
154 151 155 particle size, both the enhanced bulk modulus (γ -Fe2O3, Au, and Ag ), and the
156 157 158 decreased bulk modulus (CdSe, ZnS and TiO2 ) have been reported. In addition, PbS nanoparticles have been reported to show a bimodal dependence of particle size, as the bulk modulus value increases with the particle size up to 7 nm, and then decreases with further increasing the size of the particle.153 Therefore, there is still no agreement on how the size of the particle affects the value of the bulk modulus, and more research on other types of nanoparticles are still needed.
Optoelectronic properties of nanoparticles under high pressure have also been investigated by absorption and photoluminescence spectroscopy. It has been reported that the surface plasmon resonance of Ag nanoparticles varies with the distance between them at mesoscale.144, 159 In addition, the bandgap of a material can be shifted significantly by the application of pressure. For example, PbS displays a red shift while CdSe displays a blue shift when pressure increases, and their bands usually broaden under high pressure.130
Also, direct bandgap WZ and zinc blend (ZB) phases of CdSe nanoparticles exhibit excitonic features in their absorption spectra and strong fluorescence, but no absorption or fluorescence peak is observed in their indirect bandgap RS phase.131, 133 These results
18 demonstrate that both the atomic scale and mesoscale structural transformation under high pressure can influence their optical properties. Furthermore, despite the limited reports, increase of electrical conductivity has been generally observed for semiconductor nanoparticles.160
1.5.4 CdS Nanoparticles under Pressure
II-VI semiconductor nanoparticles (e.g., ZnS, CdS, CdSe, ZnO) have attracted growing scientific interests over the past few decades owing to their possible applications in modern electronic and photonic devices.161-162 Cadmium sulfide (CdS), as a typical II-
VI semiconductor with a direct wide bandgap of 2.53 eV, has been extensively studied for applications such as photocatalysis, photoresistors, light-emitting diodes and window layer of junction solar cells.163-166 To date, various methods have been developed to prepare CdS nanoparticles of different sizes and shapes.167-170 CdS can be synthesized in two types of crystalline structures at ambient conditions: the cubic ZB and the hexagonal WZ structures, where the WZ structure is more thermodynamically stable.171 Under pressure, CdS can transform into the high-pressure cubic rocksalt (RS) phase. It has been reported that the
WZ lattice in bulk CdS materials transform into RS structures when pressurized up to 2.6-
3.0 GPa,131, 172-176 which is in good agreement with the theoretical calculations reported by
Tan et al.177 In addition, it has been calculated by Corll et al. that phase transition pressures of CdS powders are very similar for WZ (2.5 GPa), ZB (3.1 GPa) and ZB-WZ mixtures
(2.8 GPa).178
High-pressure studies of CdS nanoparticles have shown that pressure induced phase transition occurs at a higher pressure compared to their bulk counterparts, and the pressure of the phase transition depends strongly on the particle size. Martin-Rodriguez et al.179
19 reported that the phase transition pressure of 5 nm CdS nanoparticles from ZB to RS was at ~6 GPa. Haase et al.180 claimed that ZB to RS phase transition of 4 nm CdS nanoparticles occurred at ~8 GPa. Experiments by Chen et al.181 indicated that the transition pressure at which Cd32S50 clusters (15 Å in diameter) transformed from WZ to RS was at ~5 GPa.
Such elevated phase transition pressure observed in nanoparticles relative to the bulk sample has been generally explained by the high surface energy of nanoparticles,156 as mentioned in the previous section. On the other hand, systematic study of how the nanosizes of the CdS particles affect this structural phase transition pressure has been rarely reported: Mishara et al.182 showed that the WZ to RS phase transition pressure decreased as the particle size increased from 10 nm to 44 nm; Nanba et al.183 announced an increase in phase transition pressure with increasing CdS particle size from 40 nm to 400 nm. Since the studies display quite controversy trends, it is of general interest to explore the high- pressure phase behavior of CdS nanoparticles of other sizes.
Besides the size effect, previous studies have also revealed that metal doping can alter the onset transition pressure for CdS nanoparticles. For example, research by Zhao et al. reveled that doping with Eu3+ increased the WZ to RS phase transition pressure of CdS nanoparticles from 4.76 GPa to 5.22 GPa;184 while doping with Co2+ reduced the ZB to RS phase transition pressure of CdS nanoparticles from 4.89 GPa to 4.06 GPa.185 While these reports with spherical CdS nanoparticles have resulted in valuable insights, nanoparticle shape has also been considered as an important parameter to tune their high-pressure properties. The effect of shape on the phase transition pressure has been reported for other
139 186 types of nanoparticles, such as CdSe and TiO2. However, no general trend of how particle shape correlates with phase transition pressure can be summarized since the
20 amount of reports is quite limited. Therefore, a systematic study to investigate how the size and shape of CdS nanoparticles affect their phase transition properties has been conducted, details will be presented in Chapter 4.
1.6 Motivations of My Projects
As mentioned above, even though lots of research have been conducted to understand the synthesis and self-assembly of various nanoparticles or nanocomposites, there is still not an established approach widely accepted to fabricate organic/inorganic hybrid nanostructures (e.g. polymer/nanoparticle nanocomposites, metal-organic framework nanostructures) with precisely controlled domain shapes and sizes. In addition, most nanomaterials synthesis methods are developed at ambient pressure conditions, since pressure induced compression can provide a unique possibility to control the structure and properties of nanomaterials without altering their chemical compositions, understanding the effect of shape and size of nanoparticles on the dependence of their high-pressure behaviors could thus provide insights into their structure-property relationships and will be potentially useful for technological developments.
The main objective of this thesis is to synthesize and self-assemble nanoparticles and nanocomposites into well-ordered nanostructures, and to understand their size- and shape- dependent high-pressure properties (Figure 1.8). Chapter 1 gives a brief overview of current methods in the synthesis, self-assembly and high-pressure studies of nanoparticles and nanocomposite materials. Each subsequent chapter will focus on a particular type of nanoparticle or nanocomposite material. Chapter 2 of this thesis discusses the self-assembly of polymers and nanoparticles into precisely controlled hybrid nanofibers by using non-covalent interactions as driving forces. Chapter 3 presents the achievement
21 of forming MOF 1D nanostructures through interfacial synthesis templated by nanoporous polymer membranes. Chapter 4 focuses on the effects of size and shape on the phase transitions of CdS nanoparticles.
Figure 1. 8 Scheme illustrating the outline of this thesis.
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43 Chapter 2. Bottom-Up Approach for Precisely Nanostructuring Hybrid
Organic/Inorganic Multicomponent Composites for Organic
Photovoltaics
(Reproduced with permission from Polymers 2016, 8, 408, Copyright © 2016, MDPI,
Basel, Switzerland; Nanoscale Adv. 2020, 2, 2462-2470, Copyright © The Royal Society
of Chemistry 2020; MRS Adv. 2020, just accepted, Copyright © Materials Research
Society 2020
The other coauthors, Dr. Brad W. Watson II, Chris Fetrow, Dr. Hongyou Fan, Dr.
J. Matthew D. Lane are acknowledged.)
2.1 Introduction
Research and development on π-conjugated polymers (CPs) in organic optoelectronic devices have grown rapidly owing to their mechanical flexibility and their potential for large-scale and low-cost production. So far, CPs have been successfully used as key active materials in sensors,1 organic film effect transistors (OFETs),2-3 organic photovoltaic cells (OPVs),4-6 supercapacitors7 and organic light emitting diodes
(OLEDs).8-10 Among various CPs, poly(3-hexylthiophene) (P3HT) is a model material and has been extensively studied because of their excellent optoelectronic and mechanical properties, high crystallinity, high carrier mobility and solution processability. To date,
P3HT has become the standard donor material for OPVs, and exciting progress has been made to combine P3HT with soluble derivatives of fullerene, such as [6,6]-phenyl-C61- butyric acid methyl ester (PCBM) to fabricate bulk heterojunction (BHJ) devices, which have potential to be made lightweight and flexibile.11 However, P3HT-based OPVs are still
44 less mature technologies and far from commercialization due to their poor device performances compared to inorganic semiconductor-based solar cells. Therefore, tremendous efforts have been devoted to understanding and improving the device performances.
In general, solution processed CP thin films display semi-crystalline nature which makes them highly disordered.12-13 Such structural disorder could lead to an inefficient charge transport with low carrier transport mobilities. Accordingly, the formation of highly ordered P3HT solid state has been considered as an essential research aim for the development of organic optoelectronic devices. Various solution treatments and film deposition strategies have been developed and employed to induce the alignment of P3HT and in turn control their thin film morphologies.14-16 One promising method takes advantage of the fact that P3HT and P3HT-based block copolymers (BCPs) can self- assemble through solution crystallization into one dimensional (1D) nanofiber (NF) or nanowhisker structures because of the π-π interactions between their rigid thiophene backbones.17-32 These ordered 1D morphologies have been proven to possess high charge- carrier mobilities and improved OPV33-35 and OFET36-40 device performances have also been reported.
On the other hand, combining semiconductor nanocrystals with CPs represent one of the most promising strategies for long-term developments of optoelectronic devices as it can combine the advantages of both classes of materials into a single entity, and can obtain new properties that can act collectively and can be enhanced relative to that of individual building blocks. To date, nanocomposites of P3HT and inorganic semiconductor nanocrystals have been prepared by either placing P3HT in direct contact with functional
45 inorganic nanoparticles via chemical coupling or by hybridizing P3HT in a controllable manner with added nanoparticles.41-43 Most of the reported hybrid P3HT-based PV cells have been realized with CdSe,44-48 and CdTe49-50 nanocrystals as electron acceptors.
However, simple mixing of P3HT with semiconductor nanocrystals often leads to uncontrolled aggregation of nanoparticles, which results in disordered active layer morphology and low device efficiencies.
Examples have been given to covalently attach fullerene derivative nanoparticles51-
54 and inorganic nanoparticles55-56 onto CP backbones to further achieve control over spatial arrangements of donor and acceptor domains in PV cells. However, in these examples, the particle loading ratios are generally low due to their limited solubility and strong aggregation tendency. Therefore, we present here an alternative approach for the preparation of stable and well-ordered CP/nanoparticle blends through strong cooperation of several non-covalent interactions including BCP crystallization, nanoparticle aggregation, π-π interactions, complementary hydrogen bonding and coordination interactions. Our previous work demonstrated that through the crystallization-driven self- assembly of polythiophene based conjugated BCP, functionalized organic fullerenes could be non-covalently attached onto the self-assembled BCP NFs and improved solar cell BHJ morphology and device thermal stability were achieved.57-61
In this study, BCPs with P3HT backbone and hydroxy or pyridine as side chain functional groups were designed and synthesized. Mixed solvent approach was used to achieve the crystallization of BCPs into NFs. Different nanoparticles, including CdSe quantum dots (QDs)62 and iron oxide nanoparticles (IONPs)63 were then attached onto BCP
46 NFs through hydrogen bonding, π-π interactions, and/or coordination interactions to form well-controlled hybrid polymer/nanoparticle core/shell composite NFs.
2.2 Synthesis and Characterization
2.2.1 Synthetic Procedures
All reagents and solvents were purchased from Sigma-Aldrich, TCI America, or
Alfa Aesar, and used as received. [6,6]-Phenyl-C61-butyric acid methyl ester (PCBM) was purchased from American Dye Source. Tetrahydrofuran (THF) was dried by distillation from sodium-benzophenone before use.
Poly(3-hexylthiophene) (P3HT). A flame dried 100 mL round bottom flask equipped with stopcocks, septa and magnetic stir bar was charged with M1 (0.5 g, 1.34 mmol) and LiCl (0.032 g, 0.75 mmol), and 24 mL anhydrous THF was added into the flask by syringe at room temperature. The solution was then cooled to 0 °C, and 2 M solution of i-PrMgCl in THF (0.67 mL) was added. After stirring for 30 min, the solution was warmed back to 35 °C, and Ni(dppp)Cl2 catalyst (7.5 mg, 0.0134 mmol) suspended in 2.3 mL THF was injected and stirred for 10 min. The reaction mixture was then quenched with methanol.
The resulting polymer was purified by Soxhlet extractions using methanol, acetone, hexanes, THF and chloroform. The final product was recovered by precipitation into methanol, and vacuum dried at 50 °C for 24 h (black powder, 46 % yield). 1H NMR (300.13
MHz, CDCl3): δ (ppm) = 6.98 (Th-H), 2.80(Th-CH2), 1.71 (Th-CH2CH2), 1.40 (Th-
CH2CH2[CH2]3CH3), 0.94 (Th-CH2CH2[CH2]3CH3). SEC (CHCl3, 1 mL/min): Mn = 19.5 kDa, Mw = 23.4 kDa, PDI = 1.2.
BCP1. In a 100 mL flame dried round bottom flask, M1 (1 g, 2.69 mmol) and LiCl
(0.576 g, 1.34 mmol) were pumped overnight to remove any water and oxygen. 50 mL dry
47 THF was then added into the flask and the solution was cooled to 0 °C. Next, 1.98 mL i-
PrMgCl solution (2 M in THF) was injected into the flask via syringe, and the mixture was stirred for 30 min (solution 1). In another 25 mL flame dried round bottom flask, M2 (0.134 g, 0.268 mmol) and LiCl (0.058 g, 0.134 mmol) were added, and the flask was degassed under vacuum overnight. 5 mL anhydrous THF was added to the reaction mixture and the solution was stirred for 30 min at 0 °C. Solution 1 was then heated up to 35 °C, and
Ni(dppp)Cl2 catalyst (7.6 mg in 2.3 mL THF) was added, and the solution was stirred for
30 min. 0.3 mL aliquot was then taken and quenched into excess EtMgBr. SEC (CHCl3, 1 mL/min): Mn = 37.9 kDa, Mw = 42.7 kDa, PDI = 1.1. Solution 2 was then transferred into solution 1 via cannula transfer. After 45 min, the reaction was quenched by adding 2 mL
EtMgCl (2 M in THF). The polymer was precipitated into methanol, and purified by
Soxhlet extractions with methanol, acetone, hexanes, THF and chloroform. The final product was then precipitated into methanol, collected by filtration, dried overnight as a black powder (56% yield). 1H NMR (300.13 MHz, CDCl3): δ (ppm) = 6.98 (Th-H), 3.65
(-O-CH2CH2-), 2.80 (Th-CH2), 1.71-0.83 (alkyl-H’s). SEC (CHCl3, 1 mL/min): Mn = 46.6 kDa, Mw = 52.7 kDa, PDI = 1.1.
BCP2. In a dry, 50 mL round bottom flask, 150 mg BCP1 was dissolved in 120 mL dry THF and stirred under nitrogen at 60 °C for 30 min. Afterwards, tetrabutylammonium fluoride (TBAF) solution (1.1 mL, 1M in THF) was added dropwise via syringe and the solution was stirred at 60 °C for 9 h. The polymer was recovered by precipitation into methanol and then dried overnight under vacuum (black powder, 86% yield). 1H NMR
(300.13 MHz, CDCl3): δ (ppm) = 6.98 (Th-H), 3.66 (CH2OH), 2.80 (Th-CH2), 1.71-0.83
(alkyl-H’s). SEC (CHCl3, 1 mL/min): Mn = 32.9 kDa, Mw = 37.8 kDa, PDI = 1.2.
48 BCP3. In a 100 mL Schlenk flask, 65.1 mg BCP2, 23 mg 4-dimethylaminopyridine
(0.19 mmol) and 15 mL anhydrous chlorobenzene were mixed and heated to 90 °C and stirred for 30 min. Nicotinoyl chloride hydrochloride complex (17.1 mg, 0.09 mmol) was then added as a solid, and the solution was stirred at 90 °C for 8 h. Finally, the crude polymer product was precipitated into methanol, and then purified by sequential Soxhlet extractions with methanol, acetone, hexanes, THF and chloroform. The final product was isolated from the chloroform extraction, precipitated into methanol, and dried at 50 °C under vacuum for 24 h (black powder, 91% yield). 1H NMR (300.13 MHz, CDCl3): δ
(ppm) = 9.27, 8.75, 8.28 (Py-H’s), 6.98 (Th-H), 4.37 (-CH2OOC-), 2.80 (Th-CH2), 1.71-
0.83 (alkyl-H’s). SEC (CHCl3, 1 mL/min): Mn = 44.2 kDa, Mw = 51.8 kDa, PDI = 1.2.
CdSe Quantum Dots. CdSe quantum dots were synthesized by using modified procedures from previous reports.64 Selenium precursor was prepared by mixing Se powder
(0.518 g, 6.56 mmol) and tributylphosphine (1.62 g, 8.01 mmol) in a scintillation vial for
30 min. In a 50 mL three-neck round bottom flask, CdO (0.042g, 0.33 mmol), stearic acid
(0.386 g, 1,36 mmol) hexadecylamine (3.88 g, 16.07 mmol), and trioctylphosphine oxide
(3.88 g, 10.04 mmol) were mixed and heated with stirring up to 150 °C under flowing nitrogen until the initial reddish-brown solution became optically clear. Next, the reaction solution was heated to 320 °C, and selenium precursor was quickly injected into the reaction flask. Upon injection, the solution temperature dropped to 290 °C. After 2 min, the reaction flask was cooled down to room temperature by removing the heating mental.
CdSe quantum dots were recovered by precipitation with acetone, and then washed three times with hexane/acetone mixture. The final product (yellowish powder) was vacuum dried overnight and re-dispersed in hexane.
49 Phenyldithiocarbamate (PDTC) Ligand. A 50 mL round bottom flask was charged with concentrated ammonium hydroxide (30 mL, 0.435 mol) and a stir bar under flowing nitrogen. Carbon disulfide (5 mL, 0.055 mol) was then added dropwise by syringe.
Next, 10 mL ethanol was added into the reaction flask. The solution was then immersed in an ice bath and aniline (5 mL, 0.083 mol) was added dropwise over 5 min. After 45 min, the reaction mixture was warmed back to room temperature. The solvent was vacuum dried, and the remaining solid was washed with chloroform. The final yellow/white powder was vacuumed dried and stored in the refrigerator (85 % yield). 1H NMR (300.13 MHz,
CDCl3): δ (ppm) = 7.49-7.32 (Ph-H’s), 7.29-7.26 (-NH-).
CdSe Quantum Dots Ligand Exchange. 52 mg CdSe quantum dots, 4 mL dichloromethane and a few drops of hexane were added to a 20 mL scintillation vial. After completely dissolving, the solution was injected into another 20 mL vial containing PDTC ligand (4.59 g, 2.68 mmol). The mixture was then stirred at room temperature for 82 h in the dark. After the reaction, CdSe quantum dots were recovered by precipitation into methanol, and then washed three times with hexanes, followed by centrifugation. The final light yellow powder was vacuum dried and stored in the glovebox (33% yield).
Iron Oxide Nanoparticles (INOPs). IONPs were synthesized by using modified procedures from previous reports.65-66 8 nm IONP-OA were synthesized by mixing 161.5 mg (0.46 mmol) Fe(acac)3 with 2.65 mL oleic acid, 3.2 mL oleylamine and 12 mL 1- octadecene in a three-neck flask. The mixed solution was heated to 110 °C and kept under vacuum for 30 minutes. Then the mixture was heated to 295 °C and kept for 1 hour. After the reaction, the solution was naturally cooled down to room temperature and a mixture of hexane, ethanol and isopropanol was used to precipitate the NPs. The NPs were then
50 separated by centrifugation and washed three times. Finally, IONP-OA were dried in vacuum and re-dispersed in hexane. By increasing the amount of metal acetylacetonate precursor, 20 nm IONP-L-OA nanoparticles were obtained.
IONPs Ligand Exchange The synthesis of citric acid coated IONP-CA was conducted according to previously published procedures.67 120 mg IONP-OA were dispersed in 15 mL of 50/50 mixture of dichlorobenzene and N,N-dimethylformamide.
Next, 0.1 g citric acid was added, and the mixture was stirred at 100 °C for 24 hours. It was later allowed to cool down to room temperature. The NPs were then precipitated by adding ethyl ether, and then separated via centrifugation, following by washing with ethyl ether for three times.
Preparation of Hybrid Nanofibers. P3HT and BCP NFs were fabricated through a mixed solvent approach. Typical procedures were as follows: 5 mg polymer was first dissolved in 0.4 mL chlorobenzene, before 0.1 mL acetone was added slowly with stirring.
The mixture was further stirred at room temperature for 9 hours. Next, CdSe QDs or IONPs were added into the as-formed polymer NF solution with predetermined polymer/nanoparticle weight ratios and stirred for 1 hour. The resulting hybrid NF solutions were directly used for solar cell fabrication and diluted 100 times using solvent mixtures of chlorobenzene and acetone (4/1, vol./vol.) and drop-cast onto carbon coated grids for TEM analyses.
2.2.2 Characterizations
All NMR spectra were recorded on a Bruker Avance III 300 MHz spectrometer and referenced internally to the residual solvent signals. Size exclusion chromatography (SEC) was performed on a Waters 1515 system equipped with a 2414 refractive index detector
51 and a 2707 auto-sampler. The mobile phase was chloroform with 0.5% (v/v) triethylamine passing through two styragel columns (Polymer Laboratories, 5 μm Mix-C) at a flow of 1 mL/min, kept in a column heater at 35 °C. SEC results were calibrated by external polystyrene standards (Varian). Ultraviolet-visible (UV-Vis) absorption spectra were recorded on a Shimadzu UV-2401 PX spectrometer over a range of 300-900 nm using quartz cuvettes. Fluorescence emission spectra were measured using a Varian Cary Eclipse fluorimeter. X-ray diffraction (XRD) patterns were recorded using a Rigaku SmartLab diffractometer. The FT-IR spectrum was obtained using a Thermo Nicolet 380 FTIR spectrometer with a powder sample in the ATR mode. Transmission electron microscopy
(TEM) images and selected are electron diffraction pattern were taken by JEOL-202 microscope operating at 200 kV. Samples were prepared by drop casting diluted sample solutions onto a carbon coated copper grids.
2.2.3 Solar Cell Fabrication and Measurement
Indium-tin-oxide (ITO) coated glass substrates (China Shenzhen Southern Glass
Display Ltd., 8 ohms/sq) were cleaned sequentially in detergent, DI water, acetone, isopropyl alcohol (15 min each), and then treated by UV Ozone (Novascan PSD series) for
45 min. Subsequently, MoO3 (10 nm) was then deposited onto the ITO surface using an
Angstrom Engineering Amond deposition system with a vacuum level < 7 ´ 10-8 Torr.
Blend solutions were prepared by stirring predetermined weight ratios of polymers, nanoparticles, and PCBM in chlorobenzene at 100 °C for 10 h in a nitrogen glovebox. The active layers were casted from these blend solutions onto the MoO3 layer by spin coating at 500 rpm for 30 s. After that, 100 nm Al electrode was thermally evaporated through patterned shadow masks. Current-voltage (J-V) characteristics of solar cells were
52 measured by a Keithley 2400 source meter under simulated AM 1.5 G irradiation (100 mW/cm-2) provided by a Xe arc-lamp based Newport 67005 150-W solar simulator system
(Franklin, MA, USA) equipped with an AM 1.5 filter, the light intensity was calibrated by a Newport thermopile detector (model 818-010-12) equipped with a Newport 1916-C
Optical Power Meter.
2.3 Results and Discussions
2.3.1 Synthesis of Polymer Nanofibers and Inorganic Nanoparticles
OSi ) 6 2 OSi (CH ) 6 2 C6H13 C H (CH 6 13 C6H13 iPrMgCl / LiCl ClMg S Br S I Br Ni(dppp)Br S Ni(dppp)Cl * S * S S y 2 n M 2 n x m M 1 P3HT BCP 1 C6H13
H S N Bu4NF / H2O O S NH NH2 4 CS2/NH4OH C6H12O N C6H13 C H C6H12OH O 6 13 PDTC S S * S S y Cl N y n x m * S S DMAP/Chlorobenzene n x m C6H13 C6H13 BCP 3 BCP 2
Figure 2. 1 Synthetic scheme of block copolymers and CdSe QDs capping ligands.
The basic outline for the syntheses of two new BCPs are shown in Figure 2.1. The hydroxy groups in BCP2 were easily obtained through quantitative desilylation reactions of a polymer precursor and can be used as a facile synthetic handle for various functionalities, such as the pyridine moieties in BCP3 that possess high polarity and the ability to coordinate to metal-containing nanoparticles. The molecular weights of BCP2 and BCP3 are estimated by size-exclusion chromatography (SEC) to be ca. 32.9 kDa and ca. 46.6 kDa, respectively. The non-functionalized versus functionalized block length ratio (n/m, Figure 2.1) in both polymers is estimated from SEC to be ca. 4.5 to 1. Through
NMR analyses, the shorter functionalized block contains a statistical mixture of 3-
53 hexylthiophene units and functionalized thiophene units in a ca. 5/4 ratio, which leads to an overall functional group concentration of ca. 8% in both polymers. For comparison, a
P3HT homo-polymer having a molecular weight of ca. 37.9 kDa has also been prepared.
Nanofibers (NFs) of these homo- and block co-polymers were obtained through a mixed- solvent approach by dissolving ca. 5 mg of the polymers in 0.4 mL of chlorobenzene, a good solvent for both P3HT and the functionalized blocks, followed by the addition of 0.1 mL acetone, a poor solvent for P3HT but a good solvent for the hydroxy and pyridine moieties in BCP2 and BCP3, respectively. The chlorobenzene/acetone ratio of 4/1 was optimized previously to afford the most well-defined NFs with sufficient solution stability. After stirring the mixtures for ca. 9 h, UV-vis absorption spectra (Fig. 2.2) of the diluted solutions of all three polymers show clear structured profiles having λmax values at
514, 552 and 603 nm, indicating the formation of ordered aggregates.68-69 These mixture solutions were then drop cast onto carbon coated grids, and the corresponding transmission electron microscopy (TEM) images are shown in Figure 2.3 (A)-(C). P3HT forms NFs with uniform widths of ca. 14.9 ± 1.7 nm and lengths up to a few μm. BCP2, on the other hand, forms NFs with similar average widths of ca. 15.2 ± 1.8 nm but with a large distribution of fiber lengths. We are not certain about the exact mechanisms behind this observation, and speculate that it is the relatively strong hydrogen bonding interactions between the hydroxy groups in BCP2 and acetone, which limits the formation of large polymer crystallites, i.e., long fibers. Based on the same argument, the pyridine moieties in BCP3 do not form hydrogen bonds with acetone but possess stronger dipole–dipole interactions with acetone than those from pure P3HT, which leads to the formation of NFs with intermediate lengths and widths of ca. 14.5 ± 1.5 nm. To be noted, the mixed-solvent approach does lead to less
54 uniform and sometimes ill-defined P3HT NFs than those obtained from the so-called whisker method using a single marginal solvent,17, 20-21, 23 but it allows for much higher polymer concentrations (e.g., 10 mg mL−1 in our case vs. less than 1 mg mL−1 in commonly applied whisker methods) for device relevant applications. Such a mixed-solvent approach also allows for the possibility to control nanostructure morphology by fine-tuning the polymer–solvent interactions through functional group and non-solvent variations.
Figure 2. 2 UV-vis absorption spectra of nanofiber solutions of (A) P3HT, (B) BCP2, and
(C) BCP3.
55
Figure 2. 3 Transmission electron microscopy (TEM) images of (A) P3HT NFs; (B) BCP2
NFs; (C) BCP3 NFs; (D) IONP-OA; (E) IONP-L-OA; and (F) IONP-CA. Inserts: histograms of corresponding NF widths and nanoparticle diameters sampled from 100 individual objects.
CdSe quantum dots (QDs) were prepared according to slightly modified literature procedures64 and the as-prepared QDs have an average diameter of 3.33 ± 0.3 nm by TEM analysis and trioctylphosphine oxide (TOPO) as the ligand shell. UV-Vis absorption profiles of the QDs give a λmax at ca. 596 nm. TEM image of these QDs are displayed in
Figure 2.4 (A), in which the QDs are uniform in sizes and more or less dispersed without significant aggregation. However, the as-synthesized TOPO ligands form a thick, non- conductive layer outside the QD, potentially limiting electronic communications between the organic CPs and QDs. We thus replaced these alkyl ligands with shorter phenyldithiocarbamate (PDTC) ones as shown in Figure 2.1. The TEM image in Figure
2.4 (B) shows the QDs with PDTC ligands, from which an average diameter of ca. 3.03 ±
56 0.15 nm. Such size reduction is expected from the shorter PDTC ligands. UV-vis absorption measurements gave a red-shift of λmax to ca. 605 nm, which has been previously attributed to QD to ligand charge transfer interactions.70 The TEM image shows clusters of
QDs and such aggregation effects are possibly caused by stronger interactions among the rigid phenyl groups in PDTC ligands.
Figure 2. 4 TEM images of (A) CdSe QDs with TOPO ligands (B) CdSe QDs with PDTC ligands.
Iron oxide nanoparticles (IONPs) were prepared through thermal decomposition of
Fe(acac)3 at high temperatures in the presence of oleic acid and oleylamine (OA) as surface stabilizing ligands. By varying the relative amount of iron precursors while keeping other reaction conditions constant, we obtained IONPs in two different sizes, namely IONP-OA and IONP-L-OA having diameters of 7.7 ± 0.9 nm and 20.3 ± 3.2 nm, respectively. Figure
2.3 (D)-(E) show the TEM images of these two IONPs, which are well dispersed without significant aggregation due to the long aliphatic chains of capping OA ligands. The selected area electron diffraction (SAED) and powder X-ray diffraction (PXRD) patterns of IONP-
OA are shown in Figure 2.5, which confirm that the IONPs as prepared are in
Fe3O4 magnetite phase. We also performed ligand exchange reactions on IONP-OA with citric acid, and the TEM image of the resulting IONP-CA from acetone solutions is shown in 2.3 (F). IONP-CA appears smaller than IONP-OA with an average diameter of 5.0 ± 0.9
57 nm, which is understandable considering the much shorter citric acid capping ligand in
IONP-CA. Significant aggregation is also observed for IONP-CA, likely caused by the strong hydrogen bonding interactions among the surface carboxylic ligands.
Figure 2. 5 Selected area electron diffraction (SAED) pattern (A) and powder X-ray diffraction spectrum (B) of IONP-OA.
2.3.2 Self-Assembly of Hybrid Conjugated Polymer/Quantum Dot Composite
Nanofibers
To study the self-assembly processes of polymer NFs with CdSe QDs, we added equal weight of QDs into the pre-formed NF solutions. The resulting NF/QD solutions were diluted and cast onto carbon-coated TEM grids and the images are shown in Figure
2.6 (A)-(B). In the case of P3HT NFs, the CdSe QDs are found to preferentially located in areas where the NFs are present. However, most of the QDs are not closely associated with the NFs. This phenomenon is understandable since the side-chains in P3HT and TOPO ligands are both alkyl chains so that weak hydrophobic interactions bring these two components near each other. On the other hand, there are no specific interactions between these two compounds so that they are not strongly associated with one another. Things are quite different when BCP3 NFs are applied since the pyridine functionalities should have stronger coordinating interactions with the inorganic QDs and we expected to observe
58 closer interactions between the NFs and QDs. Indeed, as seen in Figure 2.6 (B), the CdSe
QDs are also concentrated in areas where the NFs are present and most of the QDs are attached to the peripheries of the NFs. Such core/shell organic/inorganic composite NF structures provide a facile means to control the nanostructures and morphologies of hybrid materials. In case of the PDTC coated CdSe QDs, they are soluble in chlorobenzene but only poorly dissolved in the chlorobenzene/acetone mixture used for NF formation. For self-assembly studies, an alternative route was taken by dissolving the polymers and QD
(1/1, w/w) in chlorobenzene first and then adding acetone. The process was monitored by using UV-vis absorption spectroscopy and showed very similar behaviors compared with those without QDs. TEM images of the resulting composite solutions are shown in Figure
2.6 (C)-(D). Both P3HT and BCP3 form similar NFs as those prepared in the absence of
QDs. Very few QDs were found where most P3HT NFs reside Figure 2.6 (C), indicating no specific interactions between these two components. On the other hand, in Figure 2.6
(D), large quantities of QDs are clearly found near the BCP3 NFs and seemingly line up along both sides of the NF with similar distances. Such behaviors can only be explained by the non-covalent interactions between BCP3 and QDs bearing PDTC ligands.
59 Figure 2. 6 TEM images of (A) P3HT NFs with CdSe QDs having TOPO ligands (1/1, w/w); (B) BCP3 NFs with CdSe QDs having TOPO ligands (1/1, w/w); (C) P3HT NFs with
CdSe QDs having PDTC ligands (1/1, w/w); (D) BCP3 NFs with CdSe QDs having PDTC ligands (1/1, w/w).
2.3.3 Organic Solar Cells Fabricated from Hybrid Conjugated Polymer/Quantum
Dot Composite Nanofibers
Table 2. 1 Summary of Solar Cell Device Performance Dataa
b c 2 d e f Blends PCE (%) Jsc(mA/cm ) Voc (V) FF (%)
P3HT BHJ 0.17 ± 0.03 1.89 ± 0.16 0.27 ± 0.00 33 ± 3 (0.19) (2.00) (0.27) (31) P3HT NF 0.53 ± 0.21 3.95 ± 1.54 0.51 ± 0.02 27 ± 2 (0.79) (5.99) (0.54) (28) BCP3 BHJ 0.78 ± 0.08 4.23 ± 0.62 0.49 ± 0.01 38 ± 1 (0.87) (4.92) (0.50) (40) BCP3 NF 0.42 ± 0.06 3.26 ± 1.15 0.34 ± 0.04 41 ± 14 (0.46) (4.35) (0.39) (58) a All numbers are reported as averages from at least five devices, highest values are included in parentheses; bAll blends have polymer/QD/PCBM weight ratio at 1/1/1 and
c d thermally annealed at 150 °C for 10 min under N2; Power coversion efficiency; Short circuit current density; eOpen circuit voltage; fFill factor.
Organic solar cell devices were fabricated using the conventional device structure:
ITO glass/MoO3 (10 nm)/active layer (100 nm)/Al (80 nm). The active layers contain polymers, either P3HT or BCP3, CdSe QDs having PDTC ligands, and PCBM at a constant
1:1:1 weight ratio for better comparison. All devices were thermally annealed at 150 °C for 10 min under N2 and the results are summarized in Table 2.1. BHJ devices are simple blends of all components from a common solution in chlorobenzene, in contrast to the NF devices, in which polymer NFs were formed first in chlorobenzene/acetone mixtures before
60 QDs and PCBM were added. As summarized in Table 2.1, addition of QDs significantly decrease device performances when compared with binary devices of P3HT/PCBM we
59 reported recently. The devices suffer greatly from both open circuit voltage (VOC) and fill factors (FF) values, indicating severe energy loss during charge separation and transport processes. Weiss et al. have recently studied the ligand shell effects on electronic properties of QDs and found that PDTC ligands act as hole acceptors when combined with CdSe
QDs.71 Based on such energy landscape, these QDs may in fact act as recombination centers and significantly reduce obtained voltages while decrease diode ideality. Although the BCP3 devices performed slightly better than P3HT devices, which is likely due to better morphologies from the self-assembly behaviors, QDs with a different ligand sets that have the correct energy alignment with both P3HT and PCBM are needed to truly investigate the effectiveness of the ternary core–shell NF structures on device performance.
2.3.4 Self-Assembly and Magnetic Responses of Hybrid Conjugated
Polymer/Magnetic Nanoparticle Composite Nanofibers
The self-assembly of polymer NFs and IONPs was conducted by adding IONPs to the pre-formed polymer NF solutions with a polymer/IONP weight ratio of ca. 2/1. The solutions were then stirred at room temperature for 1 h before being diluted 100 times with chlorobenzene/acetone solvent mixtures (4/1, v/v) for TEM analyses. Representative TEM images of these hybrid nanostructures are assembled in Figure 2.7 and table 2.2 summarizes the average numbers of nanoparticles associated with one polymer NF by sampling about 50 individual NFs.
Table 2. 2 Average numbers (Navg) and maximum numbers (Nmax) of IONPs closely associated with one polymer NF, from sampling ca. 50 individual NFs in TEM images
61 Navg of IONP per NF Nmax of IONP per NF P3HT NF / IONP-OA 1.1 ± 1.1 4 P3HT NF / IONP-L-OA 0.3 ± 0.5 3 P3HT NF / IONP-CA N/A N/A BCP2 NF / IONP-OA 4.9 ± 2.6 17 BCP2 NF / IONP-L-OA 3.6 ± 1.3 10 BCP2 NF / IONP-CA 3.9 ± 3.3 15 BCP3 NF / IONP-OA Unable to count 34 BCP3 NF / IONP-L-OA 4.7 ± 1.9 14 BCP3 NF / IONP-CA 9.6 ± 8.0 48
Figure 2. 7 TEM images of nanostructures from mixtures of P3HT NFs and (A) IONP-
OA, (B) IONP-L-OA, and (C) IONP-CA; BCP2 NFs and (D) IONP-OA, (E) IONP-L-OA, and (F) IONP-CA; and BCP3 NFs and (G) IONP-OA, (H) IONP-L-OA, and (I) IONP-CA.
The mixture solutions used for TEM analyses contain polymer NFs and IONPs at a ca. 2/1 weight ratio and polymer concentrations at ca. 0.1 mg mL−1. Scale bars in all: 200 nm.
For P3HT NFs, IONP-OA and IONP-L-OA seem to well disperse within the networks of polymer NFs but without apparent association between them, as shown in
62 Figure 2.7 (A)-(B), respectively. These observations are expected since P3HT NFs have no specific, except hydrophobic, interactions with IONPs coated with long aliphatic chains.
Thus, the organic and inorganic components can disperse well among each other without showing significant association. In the case of IONP-CA, irregularly shaped aggregates of a few hundred nanometers in size are observed in Figure 2.7 (C). These aggregates appear to contain both the nanoparticles and polymers, but discrete P3HT NFs are no longer observed. IONP-CAs are known to self-aggregate (Figure 2.3 (F)), which is caused by the strong hydrogen bonding interactions among surface carboxylic groups, the hydrophilicity of which also make these nanoparticles incompatible with hydrophobic P3HT NFs. Thus, the appearances of large aggregates composed of both components are somewhat surprising, and we are currently investigating such formation mechanisms.
In the cases of BCP2 and BCP3 NFs, similar behaviors were observed with all three
IONPs as shown in Figure 2.7 (D) through Figure 2.7 (I). IONP-OAs are well dispersed within the networks of both BCP2 and BCP3 NFs, and most of the nanoparticles are found to closely associate and align along both sides of the NFs. The difference is that the density of IONP-OAs is found to be higher along BCP3 NFs, with less free, unattached nanoparticles, than for BCP2 hybrid NFs. Similarly, IONP-L-OAs are well dispersed and associated with both BCP2 and BCP3 NFs, with stronger attachment and less free particles observed for the latter. Interestingly, IONP-CAs no longer self-aggregate and are found to align with both BCP2 and BCP3 NFs. We rationalize the observations as the following.
The hydroxy groups in BCP2 can form hydrogen-bonding interactions, in addition to hydrophobic interactions from the polymer main-chain, with the OA ligands on the surfaces of IONP-OAs and IONP-L-OAs. Such additional hydrogen-bonding interactions
63 lead to closer association of IONPs with BCP2 NFs than with P3HT NFs. The apparent stronger attachment of IONP-OAs and IONP-L-OAs to BCP3 NFs is likely caused by stronger interactions between the pyridine moieties on BCP3 and IONPs. Besides hydrophobic and hydrogen-bonding interactions, pyridine groups can also coordinate to the surfaces of nanoparticles and partially replace the original ligands. To study such effects, we precipitated a well-dissolved BCP3/IONP-OA (2/1, wt/wt) solution in chlorobenzene into methanol and washed the precipitate extensively with methanol in order to remove any free OA ligands. The remaining powder could be attracted to a nearby permanent magnet, confirming the presence of IONPs, but was found to be insoluble in any solvent. This can be explained by cross-linking of polymer chains with IONPs as the cross-linkers, through pyridine coordination interactions. We also performed infrared (IR) spectroscopy on the precipitated BCP2/IONP powder as well as on BCP2 and IONP-OA individually, and the spectra are shown in Figure 2.8. The signals at ca. 1710 cm−1 and between 1400 and 1600 cm−1, characteristic of pyridine moieties are clearly observed in both BCP3 and BCP3/IONP-OA precipitates, while the signals at ca. 1631 cm−1, 1561 cm−1, and 1454 cm−1, characteristic of OA ligands diminish in the spectrum of
BCP3/IONP-OA precipitate, suggesting the replacement of the original ligands. As for
IONP-CA, the carboxylic surface ligands can form hydrogen-bonding interactions with the hydroxy and pyridine groups in BCP2 and BCP3, respectively, leading to the observed NF attachment without significant self-aggregation.
64 Figure 2. 8 Infrared (IR) spectra on powders of IONP-OA (black), BCP3 (Blue), and precipitate of BCP3/IONP-OA mixture (red).
Figure 2. 9 Photographs of solutions of composite NFs next to a permanent magnetic cube at the start time and the times when solutions became mostly clear. Durations for such processes to take place are shown above arrows (s: second; m: minute; h: hour).
It is well-known that ferromagnetic Fe3O4 can become superparamagnetic when it displays single magnetic domains as nanoparticles with sizes below 20 nm; and these
65 nanoparticles can respond and self-assemble to external magnetic fields.72-74 We thus tested the magnetic responsiveness of the polymer/IONP composite NFs by placing a permanent magnetic cube (Neodymium Magnet N42, Applied Magnetics, ca. 100 Gauss at surface) next to the hybrid NF solutions. Photos were taken at the beginning and at the time when most of the solutes were attracted to the side of the magnet and the solutions became clear, the durations of which were also recorded. The results are summarized in Figure 2.9. Since all solutions are in identical vials and contain the same concentrations of polymers and NPs,
(10 mg mL−1 and 5 mg mL−1, respectively), the different times during which the solutions become clear can be used to compare relative association strengths between different polymer NFs and IONPs. For IONP-OA and IONP-L-OA, similar trends are observed for the three polymer NFs. BCP3 NFs display the fastest clearing times of 30 seconds with
IONP-OA and 4 minutes with IONP-L-OA, while the respective times for P3HT NFs are
4 and 14 minutes. This is consistent with TEM observations and the conclusion that BCP3
NFs form the strongest interactions with the OA coated NPs. For both P3HT and BCP3
NFs, it took longer for the IONP-L-OA composite solutions to clear out. This is likely because that the same amount of IONPs by weight was used in all cases and the much larger particles lead to significantly smaller number of particles relative to the number of
NFs, so that the relative attractive forces experienced by the NFs are weaker in the cases of larger IONPs. Surprisingly, it took much longer (ca. 2.5 hours) for the solutions of BCP2
NFs complexed with both IONP-OA and IONP-L-OA nanoparticles to be cleared out.
Although TEM images have suggested that the nanoparticles are more strongly attached to
BCP2 NFs than to P3HT NFs, the NFs of BCP2 are much shorter, i.e., the number of BCP2
NFs is much higher than that of P3HT NFs under the same concentrations, leading to a
66 smaller amount of NPs attached per NF and possibly bare NFs for BCP2. Thus, the BCP2 composite NFs may experience less attractive force from the magnet, and thus it took longer for the solutions to clear out. For hydrophilic IONP-CA, no clearing out events could be observed for P3HT NF solutions. This confirms the lack of interactions between
P3HT and IONP-CA, and the seeming co-aggregates observed in the TEM image (Figure
2.7 (C)) are likely a result from the solvent evaporation process during the TEM sample preparation. For both BCP2 and BCP3 composite NFs with IONP-CA, the solutions were cleared out in 4 and 10 minutes respectively. The faster time for BCP2 suggests stronger interactions between the –OH groups and nanoparticle citric acid ligands.
Figure 2. 10 Photographs of well-dissolved solutions of P3HT (left), BCP2 (middle), and
BCP3 (right) mixed with IONP-OA in chlorobenzene (2/1 wt./wt., 10 mg/mL polymer concentration) placed next to a permanent magnet cube. In each photo, the mixture solutions sit on the left and on the right are the pure IONP-OA solutions in chlorobenzene at identical concentrations.
We also tested magnetic responsiveness of well-dissolved solutions of polymers and IONP-OA nanoparticles at the same weight ratios and concentrations in chlorobenzene as those in hybrid NF solutions. The photographs of these experiments are included in
Figure 2.10. In contrary to composite NFs, the well-dissolved solutions did not show clearing out events but displacement of solutions from the far side to the near side of the
67 magnetic cube. The height differences between these edges are ca. 2.6 mm, 3.2 mm, 3.5 mm, and 4.0 mm respectively for solutions of pure IONP-OA, P3HT/IONP-OA,
BCP2/IONP-OA, and BCP3/IONP-OA. We believe the height differences are caused by solute concentration differences or gradients between the near and far sides to the magnet; the higher the concentration differences the larger the height differences. Given that all solutions contain the same amount of IONP-OA and polymers, the larger height differences observed for solutions containing polymers than that for the pure nanoparticle solution confirm the existence of interactions and associations between the two components. It is also understood that the P3HT/IONP-OA solution displays the smallest height difference among the three polymer mixture solutions due to the relatively weak hydrophobic interactions, while the BCP3/IONP-OA system shows the largest height difference caused by the stronger hydrogen bonding and coordination interactions.
2.3.5 Organic Solar Cells Fabricated from Hybrid Conjugated Polymer/Magnetic
Nanoparticle Composite Nanofibers and the Active Layer Morphology
Table 2. 3 Organic solar cell performance parameters using P3HT and BCP3 NFs in combination with PCBM and varied amount of IONP-OA.a
b 2 c d e f P3HT NF IONP (wt.%) Jsc(mA/cm ) Voc (V) FF (%) PCE (%)
0 11.97 ± 1.74 0.54 ± 0.02 52 ± 6 3.31 ± 0.29 1 10.93 ± 1.79 0.51 ± 0.01 48 ± 4 2.67 ± 0.25 5 5.59 ± 0.69 0.40 ± 0.01 51 ± 2 1.16 ± 0.20 50 2.02 ± 0.36 0.17 ± 0.03 37 ± 3 0.12 ± 0.01 BCP3 NF 0 6.93 ± 0.96 0.59 ± 0.01 46 ± 7 1.86 ± 0.27 1 3.93 ± 0.60 0.56 ± 0.02 31 ± 1 0.65 ± 0.10 a All devices are based on the following geometries: ITO/MoO3 (10 nm)/active layer (100 nm)/Al (100 nm). Active layers are obtained by spin-coating from chlorobenzene/acetone
(4/1, vol./vol.) of polymer NFs (10 mg/mL) and PCBM (10 mg/mL) with varied amount
68 of IONP-OA. Performance parameters are calculated from at least five individual cells. bWeight percentage relative to polymer. cShort circuit current density. dOpen circuit voltage. eFill factor. fPower coversion efficiency.
We next applied our self-assembled CP/magnetic nanoparticle hybrid NFs in organic solar cells (OSCs) in combination with the commonly used electron acceptor phenyl-C61-butyric acid methyl ester (PCBM). We used IONP-OA as the superparamagnetic particles in our studies since they are compatible and show varied interactions with all three polymer NFs, and their sizes are more uniform than those of
IONP-L-OA and comparable with those applied in previous literature reports.75-78 As for the polymers, we chose P3HT and BCP3 NFs for direct comparison since devices employing BCP2 and PCBM under standard conditions showed very poor performance and are thus less suitable to conduct comparative studies on the effects of IONP incorporation. We first studied the device performance of P3HT NFs with various amounts of IONP-OA, using optimized conditions for P3HT/PCBM devices (i.e., P3HT NF/PCBM,
1/1, wt/wt, thermal annealed at 150 °C for 10 min), and the results are summarized in Table
2.3. Previous reports all concluded that by the addition of a few weight percent of
75- Fe3O4 nanoparticles, the P3HT/PCBM device efficiencies were improved by up to 50%.
76, 78 In our case, devices employing P3HT NFs and PCBM gave a power conversion efficiency (PCE) of ca. 3.31 ± 0.29%, which is characteristic of this materials combination and comparable with the above mentioned reports. However, with just 1 wt% of IONP-OA added, the device PCE drops to ca. 2.67 ± 0.25%, as a result of slight decreases in all performance parameters, i.e., short circuit current (JSC), open circuit voltage (VOC), and fill factor (FF). With the additions of 5 and 50 wt% of IONP-OA, device performance further
69 decreases to nearly non-functional cells for the latter. The same trend was observed for devices based on BCP3 NFs. Without the addition of IONP-OA, the devices using BCP3
NFs and PCBM perform somewhat worse than those using P3HT NFs, giving an average
PCE of ca. 1.86 ± 0.27%. With the addition of just 1 wt% of IONP-OA, the device PCE is reduced by ca. 65% to 0.65 ± 0.10%, with the reduction in JSC as a major contributor. We thus did not attempt to further increase the amount of IONP-OA for these devices.
Figure 2. 11 TEM images of device active layers employing (A) P3HT NF/PCBM; (B)
P3HT NF/PCBM/IONP-OA; (C) BCP3 NF/PCBM; and (D) BCP3 NF/PCBM/IONP-OA.
Scale bars in all: 200 nm.
The TEM images of active layers of OSC devices employing P3HT and BCP3 NFs with 0 and 1 wt% IONP-OA are shown in Figure 2.11. Without IONP-OA, both the P3HT and BCP3 active layers show polymer NFs and bulk heterojunction (BHJ) morphologies with domain sizes on the order of tens of nanometers. With the addition of IONP-OA, no
70 large phase separation could be observed, and the nanoparticles are well dispersed in both films. So the detrimental effect from IONP-OA addition on the device performance is unlikely to be caused by significant morphological changes induced by the inorganic nanoparticles. In order to probe the microscopic packing structures of the polymers, we performed X-ray diffraction (XRD) experiments on BCP3 NF thin films with PCBM,
IONP-OA, and both, and the results are shown in Figure 2.12. All films show (100), (200) and (300) peaks from polymer lamellar stacking motifs with similar intensities. The thin films containing PCBM, with or without IONP-OA, display slightly enhanced peaks around 20° (2θ), which correspond to π–π stacking among P3HT main-chains. As a result, the device performance drop is unlikely to result from IONP induced disruption on polymer crystalline structures. Based on the above observations, we suspect that, due to the specific interactions between polymer NFs and IONP-OA, the insulating inorganic nanoparticles are strongly associated with the polymer NFs and located in between polymer NFs and
PCBM in the solid state, as suggested by TEM images, which can potentially act as a barrier for charge separation and thus decrease device performance. Confirmation of such a hypothesis will require more thorough and sophisticated photophysical studies that are currently under way and will be reported in future accounts.
71 Figure 2. 12 X-ray diffraction (XRD) profiles of thin films (100 nm in thickness, thermal annealed at 150 °C for 10 min) of BCP3 NF (black), BCP3 NF/PCBM (red), BCP3
NF/IONP (blue) and BCP3 NF/IONP/PCBM (green).
2.4 Conclusions
In summary, we have prepared two novel conjugated block copolymers, BCP2 and
BCP3, based on P3HT backbone having selectively functionalized hydroxyl or pyridine moieties, which forms well-defined NFs in mixture solvents. Self-assembly of such BCP3 and BCP3 NFs with CdSe QDs or magnetic IONPs in solutions led to the formation of core–shell organic/inorganic composite NFs. Such strategy provides a unique opportunity to control the special arrangement between incompatible components and potential benefits in organic electronic devices including photovoltaics. When BCP NFs were combined with
QDs, the PDTC ligand shells of nanoparticles were found to act adversely toward OSC performances in ternary polymer/QD/PCBM devices and we are currently investigating
72 other possible ligand sets for QDs that have the correct electronic requirements to allow desired charge transfer processes. In case of incorporating IONPs, the resulting hybrid NFs display electronic properties from the polymer and magnetic responsiveness from the nanoparticles. OSCs were fabricated using these hybrid nanofiber systems, but it was found that addition of iron oxide nanoparticles detrimentally affected device performance, which is likely caused by their insulating nature that limits charge transfer efficiency. Our methodology is general and can be applied to a wide range of conjugated polymers and inorganic nanoparticles, where controlled morphologies on the nanometer scales are necessary.
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83 Chapter 3. Metal-Organic Framework (MOF) One-Dimensional
Nanostructures
(This chapter is based on a manuscript in preparation)
3.1 Introduction
Metal-organic frameworks (MOFs), as an emerging and rapidly growing class of porous materials, are assembled through metal ions or metal ion clusters linked by coordinated organic bridging ligands.1 They have attracted significant research interests in recent years because of their high internal surface area, high pore volume, and easily tunable structures, porosity and surface functionalities.2-3 Those advantages make MOFs promising candidates for a variety of applications including gas separation and storage,4-7 adsorption,8-9 catalysis,10-11 sensors,12-13 drug delivery14-15 and energy storage.16-17
Recently, much research attention has been given to nanoscale MOFs as they have potentials to share the unique properties of both porous materials and nanostructures.
Because of their high ratio of exposed active sites and rapid adsorption/desorption kinetics,
MOF nanoparticles have been used in areas such as catalysts and biomedicines.18-21 In addition, complex superstructures can be produced from the assembly of MOF nanoparticles.22-23 Such applications require the synthesis of monodispersed MOF nanoparticles with uniform sizes and shapes to provide control over the resulting porous structures on the bulk scale. However, obtaining nanoscale MOF crystallites with excellent uniformity is still a major challenge in this field as the current methods always lead to a mixture of crystals with broad distribution of size and morphology, which makes the control of the overall material properties difficult. Since MOF nanoparticles are hard to purify post-synthetically, confining the size and shape of nanoparticles during the synthesis
84 has been considered as a reliable strategy for the production of uniform nanoparticles. The commonly used synthesis methods of MOF nanostructures include rapid nucleation through fast precipitation or accelerated heating,24-25 nanoreactor confinement using emulsion systems,26-27 and coordination modulation via chemically controlling the ligand- metal interactions.28-29 As a result, various MOF nanostructures, from 0D to 3D, have all been successfully prepared.30-31
Zeolitic imidazolate frameworks (ZIFs), which are composed of imidazolate linkers and metal ions, are a novel subclass of MOFs that possess attractive properties such as crystallinity, micro-porosity, high surface area, and high thermal and chemical stability.32-
33 In particular, ZIF-8, which consists of zinc ions and 2-methylimidazolate (2-MIM) ligands, has been the research focus for gas separation applications owing to their large cavities within the framework.34-36 Besides ZIF-8, other types of ZIF structures, such as the cobalt-based ZIF-67, has also been widely researched and applied for applications such as catalysts in electrochemistry.37-38 In case of ZIF nanostructures, synthesis of 0D nanospheres and 2D membranes have both been widely reported and these materials have found use in applications such as catalysis and gas/liquid separation.39-43 However, relatively less efforts have been made to control the MOF growth in 1D direction for fabricating nanorods or nanowires. Synthetic methods toward the formation of 1D ZIF nanostructures that have been developed so far including the top-down approaches such as electrospinning,44-47 and the bottom-up approaches such as the controlled solution-based template synthesis.48-52
Another alternative strategy, which is an interfacial synthesis method based on the use of hard templates, such as porous polymers53 or porous aluminum templates54 has been
85 recently emerging as a feasible and effective synthetic route to confine MOF nanostructures with controlled morphologies. As a result, 1D MOF nanostructures can be formed within the pores of the template. Inspired by this template strategy, our group has recently reported the formation of 1D ZIF-8 nanowires, nanotubes and nanorods by using a commercially available track-etched polycarbonate (PCTE) membrane as the template.55
We found that the highly ordered cylindrical shape of the membrane pores led to a directed growth of ZIF-8 with preferential crystallographic orientations. Nanorods and nanotubes are formed in 100 nm membrane pores, and single-crystalline nanowires can be obtained within 30 nm membrane pores. This templated interfacial synthesis approach is facile and is the first demonstration of the formation of 1D MOF nanostructures by PCTE membranes.
In this study, we have tried to generalize this interfacial synthesis method by using
PCTE membranes with diverse pore sizes ranging from 10 nm to 2 µm. In addition, besides
ZIF-8, this methodology has also been applied to synthesize ZIF-67 by replacing the zinc ions with cobalt ions. X-ray diffraction (XRD), transmission electron microscopy (TEM) and scanning electron microscopy (SEM) characterizations revealed that well-dispersed
1D ZIF nanowires, nanorods, nanocylinders and nanodisks can be successfully synthesized via our methodology.
3.2 Synthesis and Characterization
3.2.1 Synthetic Procedures
Materials: Zinc nitrate hexahydrate (Zn(NO3)2•6H2O, Alfa Aesar, 99%), cobalt nitrate hexahydrate (Co(NO3)2•6H2O, Alfa Aesar, 98-102%), 1-ocatonal (Alfa Aesar, 99%), reagent grade water (BDH), and 2-Methylimidazole (2-MIM, Acros Organics, 99%) were
86 used as received without further purification. Track-etched polycarbonate membranes
(PCTE) were purchased from Sterlitech Corporation (Kent, WA) and were used as received.
Synthesis of MOF nanostructures: In a typical synthesis of ZIF-8 nanostructures, predetermined amounts of Zn(NO3)2•6H2O and 2-MIM were dissolved in reagent grade water and 1-octanol, respectively. PCTE membranes were then floated on the surface of the metal solution with the hydrophilic side (dull side) down for 24 hours. Next, 2-MIM solution was gently layered on the top of the membrane. After predetermined reaction times,
PCTE membranes were taken out, rinsed thoroughly with DI water, and dried in air. In case of synthesizing ZIF-67 nanostructures, Co(NO3)2•6H2O was used instead of
Zn(NO3)2•6H2O. The detailed reaction conditions are listed in Table 3.1.
Table 3. 1 Reactions conditions for different sizes MOF nanostructures.
Pore Size/Structure Concentration Concentration Reaction Time [Zn2+] [2-MIM] (hours) 10nm ZIF-8 0.06 1 4 30nm ZIF-8 0.025 1 1 100nm ZIF-8 0.042 2 1 200nm ZIF-8 0.025 0.5 4 2µm ZIF-8 0.1 2 4 10µm ZIF-8 0.025 1 4 20µm ZIF-8 0.025 2 4 Pore Size/Structure Concentration Concentration Reaction Time [Co2+] [2-MIM] (hours) 10nm ZIF-67 0.042 2 4 30nm ZIF-67 0.025 0.75 1 100nm ZIF-67 0.06 2 1 200nm ZIF-67 0.06 2 4 2µm ZIF-67 0.1 2 4 10µm ZIF-67 0.06 1 6 20µm ZIF-67 0.1 2 6
3.2.2 Characterizations
Transmission electron microcopy (TEM) samples were prepared by dissolving as- synthesized PCTE membranes in 8 mL chloroform, and then drop-casting on to carbon
87 coated copper grids (TED Pella Inc.). TEM images and selected area electron diffraction
(SAED) patterns were taken on a JEOL 2010F TEM at an acceleration voltage of 200 kV.
X-ray diffraction (XRD) patterns were recorded at room temperature using a Rigaku
Smartlab diffractometer with a Cu K a beam (l = 1.54 Å) operated at 40 kV and 40 mA.
SEM samples were prepared by cutting PCTE membranes in half, and then mounted onto a glass substrate using double-sided carbon tape. A layer of aluminum (about 10 nm thick) was then deposited on top of the membrane using the angstrom Engineering Amond deposition system. Isolated SEM samples were prepared by coating the hydrophobic side of the membrane with 10 nm aluminum, and then gluing the hydrophobic side onto a glass substrate using J-B Weld. Resulting samples were then soaked in THF for 10 minutes, and then taken out, dried in air. Finally, 10 nm of aluminum was deposited on the substrate.
Scanning electron microscopy (SEM) images were taken on a FET Quanta 3D FEG
SEM/FIB instrument.
3.3 Results and Discussions
Table 3. 2 Pore size and thickness of the PCTE membrane.
Pore Size Thickness (µm) 10 nm 6 30 nm 6 100 nm 6 200 nm 10 2 µm 10 10 µm 10 20 µm 3
88 The track-etched polycarbonate (PCTE) membranes with different sizes were purchased from the Sterlitech Corp., Table 3.2 summarizes the pore size and thickness of the corresponding membrane. For a typical synthesis, the membrane was put in between two solutions with the predetermined concentrations, the bottom was an aqueous solution containing metal ions (Zn for ZIF-8 and Co for ZIF-67), and the top was the 1-octanol solution containing 2-methylimidazole (2-MIM) ligands (Figure 3.1). The two precursor solutions can then diffuse into the PCTE template pores and react, leading to the nucleation and growth of corresponding 1D ZIF nanostructures within the pore channels of the membrane. The membrane was left to react for a certain amount of time before being removed from the reaction medium. Varying the experimental conditions (pore sizes, concentrations, reaction time, metal ion source, etc.) leads to different 1D ZIF nano/micro- structures with the dimensions corresponding to the pore size and thickness of the PCTE templates. The detailed reaction conditions are shown in Table 3.1.
Figure 3. 1 Image of the reaction vial for the synthesis of ZIF-8 1D nanostructures, the dotted circle represents the PCTE membrane.
89 After the synthesis, the as-prepared membrane was taken out, washed with deionized water and dried in air. To confirm the formation of ZIF-8 or ZIR-67 structures within pore channels of the membrane, we first examined the crystal structure by X-ray diffraction (XRD). Figure 3.2 shows the resulting XRD patterns. As can be seen from
Figure 3.2, the peaks of the as prepared ZIF-8 and ZIF-67 samples are in agreement with those in the corresponding simulated XRD patterns, indicating the presence of ZIF-8 and
ZIF-67 crystallites within the membrane.
Figure 3. 2 XRD patterns of as-synthesized (A) ZIF-8 and (B) ZIF-67 membranes with different pore sizes.
90 It should be noted that the relative intensities of the (011) and (002) diffractions peaks are different than the simulated pattern. To further investigate the preferential growth of the as-synthesized ZIF nanostructures, the crystallographic preferred orientation (CPO) indices of the (002) reflection in relationship to the (011) (CPO(002)/(011)) and (112)
(CPO(002)/(112)) reflections were calculated from the following equation:
CPO(002)/(011) = [(I(002)/I(011))sample - (I(002)/I(011))standard]/(I(002)/I(011))standard. (1)
The calculated CPO indices of all as-synthesized MOF nanostructures are shown in Table
3.3. The CPO indices of ZIF-8 represent a preferred orientation along the {100} plane, parallel to the porous walls of the PCTE membrane. The oriented growth of ZIF-8 crystals along the {100} crystal plane has been previously demonstrated, and the CPO values are generally reported to be higher than 50 for a strong oriented crystalline structure.56-58
However, contrary to the previously reported data, our calculated CPO values are significantly lower. This could be possibly caused by the misalignment of the X-ray beam.
Since the track-etched pores are randomly orientated, it’s difficult to perfectly align the X- ray to the crystallites formed within those pores. In case of ZIF-67, the even lower value of the COP indices are obtained, suggesting a random orientation of the as-formed nanocrystals. Reasons that caused this different crystal orientations between ZIF-8 and
ZIF-67 membranes are still under investigation. We’ve found from surface characterization of PCTE membranes by scanning electron microscopy (SEM) that even after purification, the presence of some randomly distributed crystallites at the surface of the membrane is unavoidable, especially for ZIF-67 containing membranes, which could be a possible reason for the observed randomly oriented membrane.
Table 3. 3 CPO indices for different sizes MOF nanostructures.
91 Structure CPO (002)/(011) CPO (002)/(112) 10nm ZIF-8 3.29 1.02 30nm ZIF-8 4.13 1.05 100nm ZIF-8 4.17 1.50 200nm ZIF-8 4.35 1.19 2µm ZIF-8 3.19 1.14 10µm ZIF-8 5.43 1.05 20µm ZIF-8 6.34 1.29 Structure CPO (002)/(011) CPO (002)/(112) 10nm ZIF-67 3.85 0.06 30nm ZIF-67 2.55 0.03 100nm ZIF-67 2.93 0.28 200nm ZIF-67 3.67 0.41 2µm ZIF-67 2.73 0.18 10µm ZIF-67 4.05 0.49 20µm ZIF-67 2.08 0.03
Next, scanning electron microscopy (SEM) was used to exam the surface morphology of both the hydrophobic and the hydrophilic surfaces of the resulting PCTE membranes, as shown in Figure 3.3. It should be noted that the as purchased 10 nm membranes do not have a uniform pore size, some of the pores are as large as 30 nm.
Comparing the morphologies of the hydrophobic and hydrophilic sides, it can be clearly observed from the SEM images that for 10 µm and 20 µm membranes, pores on both sides are filled, indicating that the nanostructures are formed throughout the membrane and should have the length of 10 µm and 3 µm, respectively. For membranes with pores of other sizes, the pore filling is observed at the hydrophobic side, while the hydrophilic side is not filled or partially filled, suggesting that the length of the resulting 1D nanostructures should be shorter than the thickness of the corresponding membrane. In addition, as can been seen from Figure 3.3, both the hydrophobic and hydrophilic sides contain surface materials, which cannot be completely removed without damaging the PCTE membrane.
92 Figure 3. 3 SEM images of both the hydrophobic and hydrophilic surfaces of ZIF-8 and
ZIF-67 containing PCTE membranes with different pore size. The scales of each image are shown on the left.
93
94 Figure 3. 4 (a)-(j) TEM images of ZIF-8 and ZIF-67 nanowires and nanorods formed within the pores of the PCTE membranes. Insets in (a)-(h) are TEM images with lower magnification, insets in (i) and (j) are optical microscopic images; (k)-(n) SEM images of
ZIF-8 and ZIF-67 nanocylinders and nanodisks formed within the pores of the PCTE membranes. Insets in (k)-(n) are optical microscopic images.
95
96 Figure 3. 5 SEM images of isolated nanowires, nanorods, nanocylinders and nanodisks.
The scales of each image are shown on the left.
To further characterize the morphology of the 1D nanostructures formed within the pore channels of the PCTE membrane, the membranes were then dissolved in chloroform, and SEM, transmission electron microscopy (TEM) and optical microscopy samples were prepared. Representative images of each sample are shown in Figure 3.4. A set of SEM images of as synthesized 1D nanostructures was also taken and shown in Figure 3.5. As can be seen from Figure 3.4 and Figure 3.5, 1D ZIF-8 and ZIF-67 nanowires, nanorods, nanocylinders and nanodisks have all been successfully synthesized with uniform sizes and shapes. The diameters and lengths of different types nanostructures are summarized in
97 Table 3.4. The diameter of each sample is similar to the template pores. The length of nanocylinders and nanodisks are around 10 µm and 3 µm, respectively, consistent with the template thickness. The average lengths of nanowires and nanorods are shorter than the thickness of the corresponding membrane, which are in agreement with the pore filling analysis conducted by SEM (Figure 3.3).
Table 3. 4 Diameters and lengths of MOF nanostructures.
Structure Diameter Length 10nm ZIF-8 23.92 ± 5.96 nm 1.74 ± 0.26 µm 30nm ZIF-8 28.99 ± 3.78 nm 1.65 ± 0.69 µm 100nm ZIF-8 121.83 ± 18.60 nm 2.39 ± 0.91 µm 200nm ZIF-8 225.67 ± 23.14 nm 2.60 ± 0.66 µm 2µm ZIF-8 2.12 ± 0.24 µm 7.56 ± 0.59 µm 10µm ZIF-8 7.24 ± 1.07 µm 10.67 ± 2.48 µm 20µm ZIF-8 19.22 ± 1.45 µm 2.38 ± 0.60 µm 10nm ZIF-67 23.58 ± 5.48 nm 1.78 ± 0.70 nm 30nm ZIF-67 30.54 ± 3.94 nm 1.38 ± 0.39 nm 100nm ZIF-67 71.31 ± 5.99 nm 1.98 ± 0.37 µm 200nm ZIF-67 236.92 ± 22.19 µm 4.02 ± 0.63 µm 2µm ZIF-67 1.81 ± 0.11 µm 6.92 ± 0.48 µm 10µm ZIF-67 8.20 ± 0.79 µm 9.63 ± 2.37 µm 20µm ZIF-67 18.94 ± 1.35 µm 2.88 ± 0.52 µm
98 Figure 3. 6 TEM images and the corresponding SAED patterns of (A) (C) ZIF-67 100 nm nanorods and (B) (D) ZIF-67 30 nm nanorods. Samples were calibrated with silicon standard. The red circle shows where the SAED pattern was obtained.
We have previously reported that the ZIF-8 30 nm nanowires formed from this methodology are single crystalline, while ZIF-8 100 nm nanorods are polycrystalline.55 For comparison, selected area electron diffraction (SAED) patterns were also collected for ZIF-
67 30 nm nanowires and 100 nm nanorods (Figure 3.6). As can be seen from Figure 3.6,
100 nm ZIF-67 pattern shows a polycrystalline feature. In case of the 30 nm nanowires, the SAED pattern display a single-crystalline a single-crystalline pattern along [111] zone axis, indicating that the major exposed crystal plane of the obtained nanorods is {111}. It should be noted that SAED pattern of MOF structure is difficult to obtain because of their electron beam-sensitive properties. Close examinations of the SAED patterns of other sizes nanorods and nanowires are currently underway.
The formation mechanisms of the resulting MOF nanostructures are detailed in our previous publication.55 Briefly, before adding the organic linker solution on top of the
PCTE membrane, the membrane was soaked on top of the metal-containing aqueous solution overnight. Therefore, the metal ion concentration throughout the membrane pores was considered constant from the hydrophilic side to the hydrophobic side. Once the organic solution was added on top of the membrane, the concentration of 2-MIM ligands was the highest at the hydrophobic side, and the lowest at the hydrophilic side. Therefore, the reaction rate was faster at the hydrophobic side than at the hydrophilic side, which explains why most of the hydrophobic pores are filled. In addition, the poly(N-vinyl- pyrrolidinone) (PVP) coating within the pores of the PCTE membrane was acting as the
99 anchor sites for the initially formed ZIF-8 or ZIF-67 seeds, thus the nanostructures were growing from the pore walls and proceeded inwardly. For nanowires with smaller diameters, the initially formed seeds were unstable, through an Ostwald ripening process, large single crystals were then formed.
3.4 Conclusions
In summary, by varying the pore sizes and the thickness of the purchased PCTE membranes, we have successfully prepared ZIF-8 nanowires, nanorods, nanocylinders and nanodiskers. By replacing the zinc ions with cobalt ions, we have also synthesized similar
1D nanostructures for ZIF-67. Our results have further demonstrated the generality of using this interfacial synthesis methodology templated by PCTE membranes to create MOF nano/micro- structures.
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107 Chapter 4. Size and Shape Dependence of Pressure Induced Phase
Transition in CdS Semiconductor Nanocrystals
(Reproduced with permission from J. Am. Chem. Soc. 2020, 142, 6505-6510, Copyright
© 2020 American Chemical Society; MRS Adv. 2020, just accepted, Copyright ©
Materials Research Society 2020
The other coauthors, Dr. Hongyou Fan, Dr. J. Matthew D. Lane, Luke Baca,
Jackie Tafoya, Dr. Tommy Ao, Dr. Brian Stoltzfus, Dr. Marcus Knudson, Dr. Dane
Morgan, Dr. Kevin Austin, Dr. Changyong Park, Dr. Paul Chow, Dr. Yuming Xiao, Dr.
Ruipeng Li are acknowledged.)
4.1 Introduction
Over the past few decades, numerous studies have been devoted to the design, synthesis and development of inorganic nanomaterials because of their fascinating physical and chemical properties that cannot be achieved in the bulk state.1 For example, inorganic semiconductors exhibit novel optical and electronic properties when their sizes are reduced to the nanoscale, where fundamental physical properties become size and shape dependent due to the quantum confinement of carriers and an increase in the number of surface atoms.2
In particular, II-VI semiconductor nanomaterials have gained much research attention owing to their tunable band structures and high optical absorption and emission coefficients.
This remarkable optoelectronic nature privileges them for a wide variety of potential applications, such as solar cells,3-4 light-emitting diodes and laser diodes,5-6 and biological labels.7
As one of the most studied and prepared II-VI semiconductor materials, cadmium sulfide (CdS) has fascinated generations of researchers because of its size- and shape-
108 dependent optical and electronic characteristics and polymorphous transformations between structural phases. Having a direct bandgap of 2.4 eV, CdS belongs to the class of wide bandgap semiconductors, and has been applied in many applications such as visible light sensor,8 photoresistors,9 photocatalysis10 and window layers of junction solar cells.11-
12 The natural CdS has been used as a pigment for hundreds of years because of its color in yellow and advanced thermal stability. CdS exists, like most of the group II-VI binary compounds, in three crystal forms: the hexagonal wurtzite (WZ) structure, which is the most stable one at room pressure and temperature, cubic zinc blende (ZB) structure, and the cubic rocksalt (RS) structure at high pressure.13
Extensive research during the past 20 years has led to the development of various synthetic routes that yield high quality semiconductor nanocrystals with unusual properties and structures. So far, major efforts have been focused on chemically manipulating their size, shape, composition and surface chemistry. Reaction parameters, such as temperature, time, PH, molar ratio and solvents have all been reported to play a significant role in controlling the stoichiometry, crystallinity, phase purity, size, and shape of semiconductor nanocrystals.14-15 In parallel with temperature, pressure, as another fundamental thermodynamic parameter, has also been proven to be a powerful method to induce dramatic changes in lattices parameters and electronic configurations of nanomaterials by physically adjusting their interatomic distances. The solid-solid phase transitions under pressure have been investigated in detail for many nanocrystalline materials, such as
16 17 18 19 20 21 CdSe, PbS, ZnO, TiO2, Mn3O4, etc. Besides the shrinkage at atomic lattice, pressurizing nanoparticles and their assemblies have also been shown to be effective in tuning their mesoscale structures. As pressure increases, interparticle separation decreases.
109 By further increasing pressure, nanoparticles located along the pressure applied direction can then contact each other and consolidate into new structures by nanoparticle coalescence.
Previous research has revealed that both metal and semiconductor spherical nanoparticles can sinter into nanowires under high pressure.22-30 Other nanostructures formed by nanoparticle sintering, such as 2D nanosheets31-32 and 3D interconnected network33 have also been reported. Such morphology control at the mesoscale under high pressure opens up new doors toward the nanostructure design and modification that are difficult to achieve by chemical synthesis at ambient conditions.
Previous studies on high-pressure phase transition behaviors of CdS nanoparticles have shown that the WZ to RS phase transition of CdS nanoparticles occurs at elevated pressure in comparison to the bulk sample34-36, which is in agreement with trends reported for other types of nanoparticles, and can be explained by the high surface energy of nanoparticles16. In addition to the nanosize effect, metal doping can also alter the phase transition properties of CdS nanoparticles. Prior research has revealed that doping with
Eu3+ can increase the WZ to RS phase transition pressure of CdS nanoparticles from 4.76
GPa to 5.22 GPa; while doping with Co2+ can reduce the ZB to RS phase transition pressure of CdS nanoparticles from 4.89 GPa to 4.06 GPa.37-38 Other than these reports, investigations of how the nano-sizes of CdS nanoparticles affect this pressure-depended phase transition has been rarely reported and opposite trends have been observed. Mishara et al.39 reported that the transition pressure decreased as the particle size increased from 10 nm to 44 nm; while Nanba et al.36 claimed an increase in phase transition pressure with increasing CdS particle size from 40 nm to 400 nm. Furthermore, the size effect on pressure-induced morphology transition is also not well understood. On the other hand,
110 how the shape of the particle influences the phase transition has been rarely scrutinized.
Lee et al. theoretically predicted that the phase transition pressure of CdSe nanorods 58 decreased with rod length.40 Park et al. studied the shape-dependent compressibility in rice- shaped and rod-shaped TiO2 nanoparticles.41 To more thoroughly understand the size and shape effects on high-pressure phase transition of nanoparticles, detailed experimental studies on different kinds of nanoparticles are still needed. Therefore, it is of general interest to explore the influence of CdS nanoparticle sizes and shapes on their high- pressure-induced properties in a systematic manner.
In this work, we have prepared spherical 7.5, 10.6, and 39.7 nm diameter CdS nanoparticles, and also nanospheres (5.3 nm in diameter), short nanorods (6.9 nm in diameter, and 20.1 nm in length), and long nanorods (2.9 nm in diameter and 34.9 nm in length) in the hexagonal WZ phase to study the size and shape effects by using in-situ high- pressure wide-angle X-ray scattering (WAXS) measurements. In addition, we have studied the change of nanoparticle morphologies by transmission electron microscopy (TEM) before and after the high-pressure experiments. Furthermore, bulk moduli of different particles in different phases are calculated for comparison.
4.2 Synthesis and Characterization
4.2.1 Synthetic Procedures
Chemicals: Cadmium oixde (CdO, Alfa Aesar, 99.998%), cadmium chloride
(CdCl2, Alfa Aesar, 99.998%), selenium powder (Se, Acros Organics, 99.5+%), sulfur powder (S, Sigma- Aldrich, 99.9%), n-octadecylphosphonic acid (ODPA, PCI
Synthesis, >99%), tri-n- octylphosphine oxide (TOPO, Acros Organics, 99%), tri-n- octylphosphine (TOP, Strem Chemicals, min. 97%), oleic acid (OA, Alfa Aesar, tech. 90%),
111 oleylamine (OLM, TCI America, >50%) and 1-octadecene (ODE, Alfa Aesar, tech.90%) were used as received without further purification.
Synthesis of Spherical CdS Nanoparticles: 7 nm and 11 nm spherical CdS nanoparticles were synthesized using the hot injection method following an established literature procedure with slight modifications.42 In a typical synthesis of 7 nm spherical
CdS nanoparticles, 0.5 mmol sulfur dissolved in ODE (3 mL) was injected into a solution of CdO (0.47 mmol) and OA (5 mmol) in ODE (13 mL) at 280 °C under N2. The reaction mixture was kept stirring for 1 minute, upon which 0.24 mmol ODPA and 2.64 mmol
CdCl2 dissolved in OLM (10 mL) was added. After stirring for another 2 minutes, the solution was cooled down to room temperature. Nanoparticles were then purified by centrifugation and washed with hexane and methanol mixture for three times, and dried under vacuum to give a yellow colored powder. By increasing the synthesis time up to 10 minutes, 11 nm spherical CdS nanoparticles were obtained.
40 nm spherical CdS nanoparticles were synthesized through hydrothermal
43 process. Briefly, 0.14 M Na2S in 400 mL water was added into 0.14 M Cd(OAc)2 in 500 mL water and kept stirring for 24h. The resulting yellow precipitants were kept in solution for another 24h, and then filtered out, redissolved in 60 mL deionized water and transferred into a Teflon lined stainless steel autoclave (200 mL), and sealed. The autoclave was heated to 200 ℃ and kept for 72h. After the reaction, nanoparticles were washed with water and ethanol for three times, and dried under high vacuum to give a yellow colored powder.
Spherical CdS nanoparticles of 5 nm in size were prepared following a previously reported hot injection method.42 Briefly, in a 25 mL three-neck round-bottom flask, 0.06 mmol ODPA, 2.64 mmol CdCl2 and 10 mL OLM were mixed, heated and stirred at 100°C
112 for 16 hours to produce the CdCl2-ODPA solution. In another three neck flask, sulfur precursor (0.5 mmol sulfur dissolved in 3 mL ODE) and cadmium precursor (0.47 mmol
CdO, 5 mmol OA dissolved in 13 mL ODE) were mixed at 100°C, and heated to 280°C with continuous magnetic stirring under nitrogen flow. After 1 minute of mixing, CdCl2-
ODPA was quickly injected, and the mixture was stirred for two more minutes. After the reaction, the solution was allowed to naturally cool to room temperature, and nanoparticles were collected by centrifugation. The crude product was then purified by repeated centrifugation using hexane and methanol, and dried under vacuum overnight. The final product was a yellow powder, which was then dispersed in toluene.
Synthesis of Rod-shaped CdS Nanoparticles: Rod-shaped CdS nanoparticles (7 nm in diameter and 20 nm in length) were synthesized based on a method reported by Joo
44 et al. Briefly, 6 mmol sulfur in 5 mL OLM was added into a mixture of 1 mmol CdCl2 and 10 mL OLM at 90°C with continuous stirring. The solution was then heated to 140°C; this temperature was maintained for 20 hours. After the reaction, nanoparticles were collected by centrifugation, further purified by washing with hexane/methanol three times, and then dried under vacuum overnight. The final yellow powder was then collected and dispersed in toluene.
Synthesis of Rod-Shaped CdSe/CdS Core/Shell Nanoparticles: CdSe/CdS core/shell nanoparticles (3 nm in diameter and 35 nm in length) were synthesized through a previously reported seeded growth method.45 CdSe seeds were synthesized by mixing 3.0 g TOPO, 0.29 g ODPA, 0.06 g CdO in a 50 mL three-neck flask, heated to 150°C and then kept under vacuum for 1 hour. Next, the reaction solution was heated to 300°C at a heating rate of 10°C/min under nitrogen flow. When all of the CdO had dissolved, 1.5 g TOP was
113 rapidly injected and the reaction mixture was heated to 350°C, upon which 0.058 g Se in
0.36 g TOP was quickly injected. The reaction was allowed to proceed for 2 minutes, and then the solution was cooled to room temperature. A hexane and methanol mixture was used to purify the product, followed by centrifugation. Finally, CdSe seeds were dried at room temperature overnight under vacuum and stored in glove box. In a typical synthesis of CdSe/CdS core/shell nanorods, 0.09 g CdO, 3.0 g TOPO, and 0.28 g ODPA were combined in a 50 mL round-bottom flask. The reaction mixture was then heated to 150°C and kept under vacuum for 1 hour. Next, under flowing nitrogen, the mixture was heated to 350°C, and 1.5 g TOP was injected after 15 min. Then, 0.12 g sulfur in 1.5 g TOP and
2 mg CdSe seeds in 1.5 g TOP were quickly injected into the reaction mixture. After 8 min of reaction, the solution was cooled to room temperature, and a toluene/methanol mixture was used to precipitate the nanoparticles, followed by centrifugation. After washing three times, the resulting yellow powder was vacuum dried and dispersed in toluene.
4.2.2 Characterizations
Transmission electron microscopy (TEM) images were taken on a JEOL-2010F microscope operating at 200 kV. Room pressure powder X-ray diffraction (XRD) patterns were measured using a Rigaku Smartlab diffractometer with a Cu Ka beam (l = 1.54 Å).
For size effect study, in situ HP-WAXS experiments were carried out at beamline 16-BMD of the Advanced Photon Source (APS) in Argonne National Lab (ANL) with X-ray wavelength of l = 0.41328 Å. For shape effect study, High-pressure WAXS measurements were acquired on beamline 16-ID-D (λ = 0.6199 Å) and 16- BM-D (λ = 0.41328 Å)46 at the Advanced Photon Source (APS), Argonne National Laboratory. A pair of diamond anvils was used to generate pressure up to 15 GPa with the flat diamond culets diameter of
114 300 µm. A rhenium gasket was pre-indented and laser drilled with a hole of 175 µm in diameter and 20 µm thick to serve as the sample chamber. A piece of ruby was also loaded into the sample chamber to monitor the sample pressure by a standard online ruby fluorescence. Neon gas was used as the pressure transmitting medium for the size effect study, and silicon oil was used as the pressure transmitting medium for the shape effect study. The exposure time was 30 s and the sample to detector distance was ~288.7 mm.
The diffraction patterns were collected on a Mar 345 image plate and integrated using the
Dioptas software.
4.3 Results and Discussions
4.3.1 Size Dependence of Pressure-Induced Phase Transition in CdS Semiconductor
Nanocrystals47
Figure 4. 1 TEM images of (a) 7.5 nm, (b) 10.6 nm, (c) 39.7 nm CdS nanoparticles and (d) corresponding room pressure powder XRD spectrum.
115 Figure 4.1(a)-(c) are the TEM images of as-synthesized CdS nanoparticle samples with average size of 7.5±0.9, 10.6±1.2, and 39.7±6.6 nm, respectively. The crystalline phase structure of each sample at room pressure was then measured by XRD, as shown in
Figure 4.1(d). The resulting XRD patterns show well resolved characteristic peaks of CdS hexagonal wurtzite phase (JCPDS card No. 75-1545).
116 Figure 4. 2 Representative synchrotron WAXS data during compression and decompression of (a) 7.5 nm, (b) 10.6 nm and (c) 39.7 nm CdS nanoparticles. r represents the releasing pressure, the black curve represents the WZ phase, the blue curve represents the RS phase, and the red curve represents a mixture of WZ and RS. Impurity peaks from gasket Re, neon gas and ruby are marked with asterisk.
Nanoparticles were then loaded into a DAC, and integrated synchrotron WAXS patterns for different samples during the compression and decompression process are displayed in Figure 4.2. It can be seen that all three samples possess the normal WZ structures at ambient pressure, and that with increasing pressure the corresponding diffraction peaks shifted to higher 2q value (lower d-spacing) as the result of the unit cell contraction. With further increasing pressure, a new phase, characterized by the appearance of new peaks that are indexed as cubic RS phase (JCPDS card No. 21-829) appeared. The onset of WZ to RS phase transition pressure was measured to be 7.60, 7.95, and 6.69 GPa for the 7.5, 10.6, 39.7 nm samples, respectively. The RS phase was maintained up to 15
GPa, and then pressure was gradually released. For 7.5 nm and 10.6 nm samples, the RS phase was preserved when the pressure was released back to ambient, which represents an irreversible phase transition process. For the 39.7 nm sample, some of the wurtzite peaks reappeared at r0 GPa (the fully decompressed state), indicating that the phase transition process is partially reversible. Compared with bulk CdS, which shows reversible WZ to
RS phase transition at about 2.6 GPa,48 CdS nanoparticles have higher phase transition pressures. 7.5 nm and 10.6 nm samples show similar WZ to RS phase transition pressure, while the phase transition pressure decreases with increasing the particle size from 10.6 nm to 39.7 nm. In the meantime, large particles tend to behave more like bulk material with
117 partial reversible phase transition processes. These results indicate that the size of the particle can significantly affect both the phase transition pressure and the reversibility of the phase transition process. Similar size-dependent phase transition behavior has also been reported for other types of nanoparticles and can be explained by the increase of surface energy with decreasing particle size.
Figure 4. 3 TEM images of (a) 7.5 nm, (b) 10.6 nm, and (c) 39.7 nm CdS nanoparticle samples after compression and decompression process.
After the high-pressure experiments, CdS nanoparticle samples were collected and dissolved in toluene for TEM analysis. Figures 4.3 (a)–(b) are the lattice-resolved high- resolution TEM images for 7.5 and 10.6 nm samples for which the lattice can be determined to be the (111) cubic RS crystal diffraction plan with d-spacing equal to 0.31 nm. The insets in Figures 4.3 (a)–(b) show the overall sample morphologies. The solubility of the 39.7 nm sample was limited; therefore, no high-resolution TEM image was obtained. Comparing the TEM images, we see a sphere to a rod-like morphology transformation of the 7.5 nm sample that is not evident in the other two samples, and this morphology transformation has not yet been observed in other high-pressure studies of CdS nanoparticles. Previous studies have shown that ordered fcc close-packed-spherical nanoparticles can be transformed to hexagonal packed nanowires under high pressure,24, 26, 49 because, by increasing pressure, nanoparticles located along the pressure applying direction can easily
118 contact one another and then consolidate into new nanostructures by nanoparticle coalescence. Our hypothesis is that even though nanoparticles are not closely packed in our experiments, some of the randomly arranged particles can still locate along the pressure applying direction. It has been demonstrated that particles with size £ 10 nm tend to sinter together to reduce their surface energy more than bigger particles,50 but whether bigger
(e.g. 39.7 nm) nanoparticles can sinter or not is still not clear. The detailed mechanism of this morphology transformation is still being studied.
119 Figure 4. 4 Dependence of unit cell volume on the applied external pressure for (a) 7.5, (b)
10.6, and (c) 39.7 nm CdS nanoparticles. Black dots represent the compression process and red dots represent the decompression process.
The unit cell volume change of different samples under pressure was calculated and summarized in Figure 4.4 From WZ to RS phase, about 17% volume decrease is calculated for all samples, close to the previously reported data.51 Bulk modulus is an important mechanical property that denotes the stiffness of the material. Bulk moduli of different samples in different phases were then determined by fitting the 2nd order Birch-Murnaghan equation of state.52-54
7/3 5/3 P=(3/2)B0[(V0/V) -(V0/V) ] (1)
In this equation, B0 and V0 are the bulk modulus and initial unit cell volume at room pressure, respectively. V0 can be calculated form the room pressure powder XRD data. The resulting bulk moduli are shown in Table 1.
Table 4. 1 Unit cell volumes and bulk moduli of CdS nanoparticles.
Sample Size Wurtzite (WZ) Rocksalt (RS) 3 3 V0 (Å ) B0 (GPa) V0 (Å ) B0 (GPa) 7.5 nm 99.92 57.49±0.93 161.10 85.14±0.59 10.6 nm 99.38 59.97±0.86 160.97 84.09±0.71 39.7 nm 97.69 78.90±5.02 161.10 84.36±0.46
Bulk moduli of nanoparticles at both the WZ and RS phases are higher when compared to the reported data of the bulk CdS material (B0 = 54.0 GPa for WZ and B0 =
68.0 GPa for RS),55 which agrees with the trends observed for other types of particles.49, 56-
57 In addition, we found that WZ nanoparticles show lower bulk modulus values than RS particles, indicating that they are more compressible. Also, for the WZ phase, the bulk modulus increases with increasing particle size, while those of RS nanoparticles remain
120 58 59 similar to each other. A similar trend has been observed for g-Fe2O3 and PbS nanoparticles, but the opposite behavior was also reported in ZnS nanoparticles.60 Our results are contrary to those in prior studies, but the particles studied previously were of different size ranges and covered by different surfactants. Therefore, there is still no agreement on how the size of the particle affects the value of the bulk modulus, and our results cannot be considered as a general trend in nanoparticles.
4.3.2 Shape Dependence of Pressure-Induced Phase Transition in CdS Semiconductor
Nanocrystals61
Figure 4. 5 Transmission electron microscopy (TEM) images of (a) spherical CdS nanoparticles, (b) short CdS nanorods, and (c) long CdSe/CdS core/shell nanorods.
We conducted high-pressure studies on CdS nanoparticles having three distinct shapes, i.e., nanospheres, short nanorods, and long nanorods, using in-situ synchrotron wide angle X-ray scattering (WAXS) measurements. Transmission electron microscopy
(TEM) has been applied to examine morphological changes of the samples before and after the compression-decompression cycle. In addition, bulk moduli of different samples in both WZ and RS phases are calculated.
Table 4. 2 Size and surface-to-volume ratio of CdS nanoparticles.
121 CdS Shape Average Size (nm) Surface Area (nm-1) Surface to volume ratio (nm-1) Sphere 5.3±0.9 86.9 1.1 Short Rods 6.9±0.9 (width) 512.4 0.7 20.1±5.1 (length) Long Rods 2.9±0.7 (width) 328.5 1.5 34.9±5.6 (length)
CdS nanoparticles were synthesized in three distinct shapes based on previously reported methods.42, 44-45 TEM was used to characterize the morphology of the as- synthesized nanoparticles, and representative images are shown in Figure 4.5. All three types of CdS nanoparticles are monodispersed in size and uniform in shape. The average particle size and surface-to-volume ratio of different nanoparticles are summarized in Table
4.2. The average size parameters were obtained by sampling at least 100 individual nanoparticles. It should be noted that the long CdSe/CdS core/shell nanorods are comparable with the other two samples in the current studies because the contribution of
CdSe core to the overall pressure-induced behaviors can be neglected due to its relatively small volume ratio.62-63
These CdS nanoparticles were then drop-cast onto Si wafers to form uniform films, and small pieces of the resulting films were scratched off and loaded into sample chambers of the Diamond anvil cells (DACs) for high-pressure experiments. The DAC was compressed quasi-hydrostatically up to 15 GPa using silicon oil as the pressure transmitting medium and WAXS experiments were performed after each pressure point was reached and stabilized. The resulting X-ray scattering patterns of different samples at different pressures are compiled in Figure 4.6 At ambient pressure before compression, the WAXS patterns of all three CdS nanoparticles can be indexed according to the hexagonal WZ crystal structure (wurtzite CdS, JCPDS card number 75-1545). With increasing pressures,
122 all WAXS peaks shifted to higher q values, corresponding to smaller d spacings resulting from shrinkage of the nanoparticle atomic lattice under applied pressures. Clear phase transitions, as indicated by appearances of new scattering peaks, were then observed at higher pressures. The onsets of such phase transitions occur at ca. 6.0 GPa for nanospheres, ca. 6.9 GPa for short nanorods, and 8.0 GPa for long nanorods. These observed new peaks correspond to the cubic RS crystal structure (cubic CdS, JCPDS card number 21-829) in all three cases, and RS structures were stable up to the highest pressure applied, i.e., 15
GPa. When the pressure was released back to ambient conditions, some of the WZ peaks reappeared in both cases of the nanorod samples (Figures 4.6 (b) and (c)), indicating a partially reversible phase transition process. On the other hand, the high-pressure RS phase is maintained at ambient pressure for the nanospheres (Figure 4.6 (a)), representing an irreversible phase transition behavior. Compared with bulk materials, WZ-to-RS phase transitions have been found to take place at higher pressures for spherical nanoparticles, which is commonly explained by the increased surface energy with reducing particle size or increasing surface to volume ratio.16 In the cases of our present studies, the nanospheres, short nanorods, and long nanorods possess surface-to-volume ratios at ca. 1.1 nm−1, 0.7 nm−1, and 1.5 nm−1, respectively. It is thus expected that the long nanorods show the highest phase transition pressure due to its highest surface-to-volume ratio. However, the short nanorods, having lower surface-to-volume ratio than that of the nanospheres, display relatively higher phase transition pressure. Furthermore, the WZ-to-RS phase transition was found to be irreversible in nanospheres, while such transitions appear to be partially reversible in both nanorods with different aspect ratios. Our results suggest that, besides considering nanoparticle surface energies, the shape of nanoparticles also plays an
123 important role in determining the pressure and reversibility of phase transitions. More precise determination and quantification of such shape-dependent phase transition effects will require more detailed and comprehensive studies on larger sets of nanoparticles with varying shapes, which is currently underway and will be reported in future accounts.
Figure 4. 6 Wide-angle X-ray scattering (WAXS) patterns under various applied pressure:
(a) CdS nanospheres, (b) short CdS nanorods and (c) long CdSe/CdS core/shell nanorods;
124 during compression and decompression. Pressures labeled with letter r are during decompression processes. The black, blue, and red curves represent the WZ, RS, and
WZ/RS mixture crystal structures, respectively. The red asterisks mark diffraction peaks from rhenium gaskets used in the anvil cells.
Figure 4. 7 TEM images of (a) CdS nanospheres, (b) short CdS nanorods, (c) long
CdSe/CdS core/shell nanorods after high pressure studies; and high-resolution TEM (HR-
TEM) images of (d) CdS nanospheres, (e) short CdS nanorods, and (f) long CdSe/CdS core/shell nanorods after high pressure studies.
After the high-pressure experiments, remaining residues from the DAC cells were dissolved in small amount of toluene and drop-cast onto TEM grids, and representative
TEM images are shown in Figure 4.7 CdS nanospheres showed insignificant size changes after compression. Interestingly, some of the nanospheres were observed to sinter into continuous wires that have width comparable to that of individual nanospheres (Figure
125 4.7a) and high resolution TEM (HR-TEM) image (Figure 4.7d) reveals that the crystal lattice belongs to the RS phase, consistent with the WAXS results. The connection between sintered nanospheres appears to be non-epitaxial since the lattice fringes do not match one another in adjacent spheres as observed in slightly zoomed out HR-TEM images (Figure
4.8). As for the nanorods, the general shapes remain unchanged as seen in Figures 4.7b,
4.7c and Figure 4.9. However, the lengths of both nanorods have become shorter and less uniform. The average length of the short CdS nanorods decreases from ca. 20.1±5.1 nm to ca. 16.3±4.5 nm, while that of the core/shell long nanorods reduces from ca. 34.9±5.6 nm to ca. 18.5±5.2 nm. Since the widths of these nanorods remain unchanged, we suspect that the observed shortening of nanorods are resulted from pressure induced breakage, which is more severe in the case of the long nanorods. HR-TEM (Figures 3d to 3f) reveals the presence of both the RS (d111 = 0.31 nm) and WZ (d100 = 0.35 nm) crystal structures, consistent with the WAXS data and confirms that the phase transitions of nanorods are partly reversible. (a) (b)
Figure 4. 8 TEM images of spherical CdS nanoparticles after compression.
126 (a) (b)
Figure 4. 9 TEM images of (a) short CdS nanorods and (b) long CdS nanorods after compression.
Evolution of the unit cell volumes as a function of pressure is shown in Figure 4.10.
It can be seen that there is ca. 17% volume reduction from WZ to RS crystal structure, which is in good agreement with previous studies.51 The volume change versus pressure data were then fitted into the second-order Birch-Murnaghan equation of state to calculate the bulk moduli of different samples,52-54
7/3 5/3 P=(3/2)B0[(V0/V) -(V0/V) ] (1) where B0 is the bulk modulus. V0 is the volume at zero applied pressure and can be calculated from the zero pressure WAXS data. The as-calculated bulk moduli of different samples at both WZ and RS phases are summarized in Table 4.3
Table 4. 3 Calculated unit cell volumes and bulk moduli of the three CdS samples.
Sample Size Wurtzite (WZ) Rocksalt (RS) 3 3 V0 (Å ) B0 (GPa) V0 (Å ) B0 (GPa) Spheres 98.36 57.89±1.36 158.72 87.97±1.72 Short Rods 98.92 66.67±1.89 160.91 85.45±1.32 Long Rods 98.45 67.69±0.75 162.11 88.58±1.05
127
128 Figure 4. 10 Pressure dependence of the unit cell volume for (a) CdS nanospheres, (b) short CdS nanorods, and (c) long CdSe/CdS core/shell nanorods. The black and red dots represent the compression and decompression process, respectively.
Materials that show higher bulk modulus values are less compressible. The CdS bulk material was reported to have a bulk modulus of 54.0 GPa for the WZ phase, and 68.0
GPa for the RS phase.55 The bulk moduli of all three samples in both WZ and RS phases are higher than that of the bulk CdS, which is in agreement with earlier studies reporting other types of nanoparticles.56-57 In addition, WZ particles are found to be more compressible than RS particles. Bulk moduli of nanoparticles in the WZ phase also shows shape-dependent features, with nanorods (high aspect ratio nanoparticles) being less compressible than spherical nanoparticles, while RS phase behaves similar for all shapes.
A similar trend has been observed for ZnO nanowires and nanobelts.48-49 But opposite
35 behavior was also observed for rice-shaped TiO2 nanoparticles. Therefore, there is still no agreement on how the shape of the particle affect the value of the bulk modulus, and more research on other types of particles is necessary to fully understand this phenomenon.
4.4 Conclusions
In summary, the impact of nanoparticle size on high-pressure-induced phase transitions were studied using spherical CdS nanoparticles having average sizes of 7.5, 10.6, and 39.7 nm. Synchrotron WAXS analysis shows unique size-dependent phase transition pressure and phase transition reversibility. It was revealed that the WZ to RS phase transition pressure increased with increasing particle size from 10.6 nm to 39.7 nm. Also, morphology transformation from sphere to rod was observed after the compression-
129 decompression cycle only for the 7.5 nm sample. Further calculations of the bulk modulus show that for WZ CdS nanoparticles, bulk modulus increases with increasing particle size.
In addition, we have employed high-pressure synchrotron WAXS to investigate the effects of particle shape on the phase transition behaviors of nanoparticles by applying CdS nanoparticles with three different shapes: CdS nanospheres, short CdS nanorods, and long
CdSe/CdS core/shell nanorods. The results show that the WZ to RS phase transition pressure and the process reversibility are both closely associated to the particles’ sizes and shapes. Spherical nanoparticles were found to possess the lowest phase transition pressure and showed sintering phenomena after the high-pressure studies. Both nanorods showed higher phase transition pressures despite the fact that the short nanorods have smaller surface-to-volume ratio than that of the nanospheres. On the other hand, both nanorods display similar bulk moduli in both WZ and RS phases, but differ significantly in phase transition pressures. Furthermore, the WZ-to-RS phase changes were found to be irreversible in nanospheres but partially reversible in both nanorods. These observations clearly demonstrate that the shape plays an important role in phase changes of nanoparticles under pressure.
With CdS as the model material, our work provides detailed information about the effects of particle size and shape on their high-pressure behavior, and it may help to initiate new approaches to designing nanoparticles for high-pressure applications.
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