University of Southern California Dissertations and Theses

Expanding the library of surface ligands for semiconductor nanocrystal synthesis and photovoltaic applications

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Expanding the library of surface ligands for semiconductor nanocrystal synthesis and photovoltaic applications

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Content EXPANDING THE LIBRARY OF SURFACE LIGANDS FOR
SEMICONDUCTOR NANOCRYSTAL SYNTHESIS AND PHOTOVOLTAIC
APPLICATIONS

by
Haipeng Lu
__________________________________________________________________
A Dissertation Presented to the
FACULTY OF THE USC GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(CHEMISTRY)

August 2017

Copyright 2017 Haipeng Lu
ii
Acknowledgements
I will begin by acknowledging my research advisor, Professor Richard L.
Brutchey, who truly made this possible. Richard’s incredible passion on science and our
common research interest brought me from China to USC and kept me motived
throughout my graduate study. His substantial guidance and strong support, both
scientifically and personally, have helped me to grow so much and get me to where I am
today. I appreciate everything Richard has taught me and done for me in the past five
years. Thank you so much, Richard.
I must also extend my gratitude to my other committee members, Professors Mark
Thompson, Noah Malmstadt, Brent Melot, and Barry Thompson. Their academic support
and guidance have inspired me in many ways and helped me to become a better
researcher. The deep discussion on the group theory and band structure in my qualifying
exam helped me to gain an unforgettable understanding, for which I will be forever
thankful to have experienced. I would also like to thank Professors Mark Thompson,
Stephen Bradforth and Oleg Prezhdo, along with their students, Dr. Andrew Bartynski,
Jimmy Joy and Dr. Zhaohui Zhou, for their outstanding work in the collaboration. I
should also thank all colleagues in the Brutchey group, both past and present, for all the
help and friendship during my academic career. I would especially like to thank Dr.
Matthew Greaney, and Prof. Dr. Federico Rabuffetti for mentoring me both as a chemist
and as an individual.
iii
Finally, I would like to express my deepest appreciation to my mother, father, and
brother, who provided me with tremendous love and support throughout this long
journey. They have always been encouraging me to pursue my true love. There is
absolutely no way that I would have accomplished any of this without them.

iv
Table of Contents

Acknowledgements ii

List of Tables viii

List of Figures

List of Schemes
ix

xix

Abstract xx

Chapter 1. Hybrid Polymer: Nanocrystal Solar Cells
1.1. Introduction
1.2. Introducting Inorganic Nanocrystals (NCs)
1.3. Hybrid Polymer: Nanocrystals Solar Cells
1.3.1. Mechanism and Theory
1.3.2. Surface Chemistry of Colloidal Semiconductor NCs
1.3.3. Tailoring the Interfacial Coupling and Energetics
Through Ligand Engineering
1.3.4. Charge Transport Dependence on NC Shapes and Film
Morphology
1.3.5. Optimizing Device Architecture
1.3.6. Harvesting NIR Light by Colloidal Lead Chalcogenide
NCs
1.4. Perspective and Future Direction
1.5. References

1
1
4
6
6
12
15

21

24
27

31
33
Chapter 2. Tandem and Triple-Junction Polymer: Nanocrystal Hybrid Solar
Cells Consisting of Identical Subcells
2.1. Abstract
2.2. Introduction
2.3. Results and Discussion
2.3.1. Single-Junction Device Optimization
2.3.2. Tandem and Triple-Junction Device
2.3.2.1. Interconnecting Layer (ICL) Optimization
2.3.2.2. Tandem P3HT: CdSe Nanocrystal Subcell
Thicknes Optimization
2.3.2.3. Triple-Junction Hybrid P3HT: CdSe Solar Cells
2.4. Experimental
2.4.1. Materials

42
42
44
45
47
47
48

53
55
55
v
2.4.2. Synthesis and Ligand Exchange of CdSe Nanocrystal
2.4.3. Characterization
2.4.4. Single-Junction Hybrid Solar Cell Fabrication
2.4.5. Tandem/Triple-Junction Hybrid Solar Cell Fabrication
2.4.6. Device Characterization
2.5. Conclusions
2.6. References
55
56
57
58
59
59
60

Chapter 3. Iodide-Passivated Colloidal PbS Nanocrystals Leading to Highly
Efficient Polymer: Nanocrystal Hybrid Solar Cells
3.1. Abstract
3.2. Introduction
3.3. Results and Discussion
3.3.1. Colloidal Iodide Ligand Exchange
3.3.2. Hybrid Solar Cells
3.3.3. Charge Separation Dynamics at The Hybrid Interface
3.4. Experimental
3.4.1. Materials
3.4.2. Synthesis of PbS Nanocrystals
3.4.3. Ligand Exchange of PbS Nanocrystals
3.4.3.1. Ligand Exchange with Lead Iodide (PbI2)
3.4.3.2. Ligand Exchange with Ammonium Iodide
(NH4I)
3.4.3.3. Ligand Exchange with N-Butylamine (BA)
3.4.4. Characterization
3.4.5. Photoluminescence Lifetime Studies
3.4.6. Transient Absorption Studies
3.4.7. Chemical Doping Studies
3.4.8. Hybrid Solar Cell Fabrication
3.4.9. Device Characterization
3.5. Conclusions
3.6. References

64
65
68
68
73
77
87
87
87
88
88
88

89
89
91
92
92
93
94
94
95

Chapter 4. Exposing the Dynamics and Energetics of N-Heterocyclic Carbene-
Nanocrystal Interface
4.1. Abstract
4.2. Introduction
4.3. Results and Discussion
4.3.1. Synthesis of NHC-Capped Ag, Ag2S, and Ag2Se NCs
4.3.2. Characterization of Surface NHC Ligands
4.3.3. Energetic NHC-NC Interface Exposed from Density
Functional Theory (DFT) Calculation

102
102
104
104
106
110

vi
4.3.4. Synthesis of NHC-Capped Cu, Cu2 xS, and Cu2 xSe NCs
4.4. Experimental
4.4.1. General Information and Characterization Techniques
4.4.2. Synthesis of Benzimidazolium Salts, NHC-AgBr, and
NHC-CuBr Complexes
4.4.3. Synthesis of NHC-Metal and NHC-Metal Chalcogenide
NCs
4.5. Conclusions
4.6. References
113
115
115
118

120

122
122

Chapter 5. Tunable Room-Temperature Synthesis of Coinage Metal
Chalcogenide Nanocrystals from N-Heterocyclic Carbene Synthons
5.1. Abstract
5.2. Introduction
5.3. Results and Discussion
5.3.1. Synthesis of NHC-Capped Ag2S Quantum Dots (QDs)
5.3.2. The Role of NHC-Ligands: Tailoring the Steric and
Electronic Properties
5.3.3. Characterization of Surface NHC Ligands
5.3.4. Synthesis of Colloidal Cu2 xS, Ag2Se and Cu2 xSe QDs
5.4. Experimental
5.4.1. Materials
5.4.2. Synthesis of Benzimidazolium Salts, Imidazolium
Salts, NHC-AgBr, and NHC-CuBr Complexes
5.4.3. Control Experiments
5.4.4. Synthesis of Metal Sulfide (Ag2S, Cu2 xS) QDs
5.4.5. Synthesis of Metal Selenide (Ag2Se, Cu2 xSe) QDs
5.4.6. Characterization
5.5. Conclusions
5.6. References

Chapter 6. Accessing Metastable Wurtzite InP Nanocrystals Via N-
Heterocyclic Carbene Stabilized Plasmonic Cu3 xP
6.1. Abstract
6.2. Introduction
6.3. Results and Discussion
6.4. Experimental
6.4.1. General Considerations
6.4.2. Synthesis and Characterization of Cu3 xP

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125
128
128
132

137
140
143
143
144

153
154
154
155
157
157

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163
163
165
172
172
174

vii
6.4.3. Synthesis of InP Via Cation Exchange Reaction
6.4.4. Synthesis of Zinc Blende InP NCs
6.5. Conclusions
6.6. References

Bibliography

174
175
175
176

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viii
List of Tables

Table 1.1: A summary of high performing hybrid polymer: nanocrystal
solar cells.
21

Table 2.1: Photovoltaic device parameters for single-junction P3HT: CdSe
BHJ hybrid solar cells with different active layer thicknesses.
46

Table 2.2: Photovoltaic device parameters for tandem P3HT: CdSe BHJ
hybrid solar cells with different active layer (AL) thicknesses.

Table 2.3: Photovoltaic device details for optimized single-junction,
tandem, and triple-junction hybrid solar cells based on a P3HT:
CdSe BHJ.
51

54

Table 3.1: Photovoltaic device parameters for hybrid Si-PCPDTBT: PbS
nanocrystal BHJ solar cells with different active layer (AL)
thicknesses

Table 3.2: Photovoltaic device raw data for hybrid Si-PCPDTBT: PbS
nanocrystal BHJ solar cells with different ligands

Table 5.1. Summary of Ag2S products from different NHC-AgBr
precursors.

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77

135

ix
List of Figures

Figure 1.1: Chemical structure of conjugated polymers discussed in this
chapter.
2

Figure 1.2: (a) Schematic illustration of hybrid polymer: NC BHJ solar
cells and type-Ⅱ energy level diagram, showing the dissociation
of excitons and charge transport process. (b) Calculated power
conversion efficiencies of a hybrid polymer: CdSe solar cell as
a function of the optical band gap of the donor material and
CdSe NC diameter.

7

Figure 1.3: Transient absorption spectra of a neat CdSe QDs film (yellow),
a neat PCPDTBT film (blue) and a hybrid PCPDTBT: CdSe
BHJ film (green) (  ex = 800 nm). The negative spectra signature
at 610 nm indicates that the quantum dot 1Se energy level is
populated by charge transfer from PCPDTBT. (b) Density of
electrons located on the NC conduction band as a function of
time, attained from the deconvolution of the 610 nm transient
absorption bleach.

10

Figure 1.4: Schematic illustration of surface ligand coordination chemistry
according to the covalent bond classification method.
13

Figure 1.5: Quantitative molecular orbital diagrams illustrating the bonding
interactions between surface Cd
2+
cations on CdSe QDs and σ-
donating ligands. (a) An undercoordinated Cd
2+
surface site
yields a ‘nonbonding’ state within the HOMO LUMO gap,
which is typically referred to as a ‘dangling bond’. (b) A strong
σ-donating ligand binds to surface Cd
2+
, leading to bonding and
antibonding states lay outside of the band gap, thus eliminating
surface traps. (c) A weak σ-donating ligand forms an
antibonding state within the HOMO LUMO gap, and creates a
surface trap statee.

15

Figure 1.6: Thermogravimetric analysis (TGA) traces (a), and (C H)
stretching region in the FT-IR spectra (b) of CdSe QDs with
their native ligand (NL), after colloidal ligand exchange with
pyridine (Py), and tert-butylthiol (tBT). The low organic
content and the absence of the high temperature mass loss

19

x
observed from TGA traces suggest a quantitative ligand
exchange by tBT ligands, which is in accordance with the
relative intensity of (C H) stretching based on FT-IR spectra.
(c) PL lifetime decay traces for films of neat CdSe QDs (  ex =
400 nm;  em = 650 nm). Lifetimes of 1.4, 3.3, and 4.4 ns were
measured for CdSe QDs with NL ligands, and after exchanged
with Py and tBT ligands, respectively.

Figure 1.7: Visible absorption spectra for (a) Py- and (b) tBT-exchanged
CdSe QDs on ITO electrodes and collected with decreasing
applied bias. All potentials reported here are referenced to NHE.
Initial bleaching is first observed at 0.54 and 0.74 V for Py- and
tBT-exchanged CdSe QDs, respectively. (c) Cyclic
voltammograms for CdSe QDs exchanged with Py and tBT,
indicating that a higher lying LUMO for the tBT-exchanged
CdSe QDs.

Figure 1.8: (a) TEM micrograph of CdSe multipods. (b) I V curves under
dark (dashed lines) and AM 1.5G 1 sun illumination (solid
lines). (c) Ideal morphology for polymer:nanocrystal BHJs
bearing some degree of vertical phase segregation, with donor
materials aggregating near the anode and acceptor materials
assembling near the cathode.

21

23

Figure 1.9: (a) EQE spectra of hybrid P3HT: CdSe and P3HT: CdSe with a
ZnO nanoparticle layer. (b) Optical intensity profile for the
device with the ZnO layer (  = 400 700 nm), showing that the
optical field intensity of BHJ is maximized. (c) Double-pass
absorption percentage of single-junction and tandem hybrid
PH3T: CdSe solar cells. (d) I V curves of hybrid P3HT: CdSe
single-junction (47 nm active layer), tandem (47/52 nm), and
triple-junction (47/47/56 nm) solar cells.

26

Figure 1.10: Hybrid Si-PCPDTBT: PbS solar cells with different ligands on
the surface of PbS QDs. (a) I V curves, and (b) EQE spectra.
(c) Femtosecond transient absorption spectra of three films
including neat Si-PCPDTBT film, hybrid Si-PCPDTBT:PbS
BHJ where PbS QDs were capped with NH4I ligands, and
hybrid Si-PCPDTBT: PbS BHJ with PbS QDs exchanged with
PbI2 ligands. Hybrid films were pumped at 920 nm and probed
at 1200 nm; the neat polymer film was pumped at 660 nm and
probed at 1200 nm.

31
xi

Figure 2.1: (a) TGA data for the as-synthesized and pyridine-exchanged
CdSe nanocrystals. (b) TEM micrograph of the as-synthesized
nanocrystals. Inset on the left is the size distribution, and inset
on the right is a high-resolution image showing the lattice
fringes of a single CdSe nanocrystal. (c) Schematic and cross-
section TEM micrograph of the tandem hybrid device
architecture: Glass/ITO/APEDOT:PSS/P3HT: CdSe(1)/ZnO/N-
PEDOT: PSS/P3HT: CdSe(2)/ZnO/Al.

45

Figure 2.2: Single-junction P3HT: CdSe nanocrystal BHJ hybrid solar cell
performance with different active layer thicknesses. (a) I −V
curves. (b) External quantum efficiency spectra. (c) Internal
quantum efficiency spectra.

46

Figure 2.3: (a) Simulated photocurrent from cell 1 in the tandem device as
a function of active layer thickness. (b) Simulated photocurrent
from cell 2 in the tandem device as a function of active layer
thickness. (c) Photocurrent for the tandem device taken from
the lowest photocurrent of the two cells. (d) Current mismatch
calculated from the difference in photocurrent between cell 1
and cell 2.

49

Figure 2.4: Dark I −V characteristic of single-junction and tandem P3HT:
CdSe hybrid solar cells.

51

Figure 2.5: Double-pass absorption percentage of single-junction and
tandem P3HT: CdSe hybrid solar.
53

Figure 2.6: I −V curves of P3HT: CdSe single-junction (47 nm active
layer), tandem (47/52 nm), and triple-junction (47/47/56 nm)
hybrid solar cells.

54

Figure 3.1: TEM micrograh of the as-synthesized PbS nanocrystals. Inset
on the right is a high-resolution image showing the lattice
fringes of a single PbS nanocrystal

Figure 3.2: Colloidal ligand exchange characterization of the PbS
nanocrystals. (a) vis−NIR absorption spectra, (b) TGA traces,
and (c) FT-IR spectra of the PbS nanocrystals with different
surface ligands (OA = oleate, BA = n-butylamine, NH4I, and
PbI2). (d) Normalized high-resolution I 3d XPS spectra for
NH4I- and PbI2-exchanged PbS nanocrystals.
69

71
xii

Figure 3.6: Photovoltaic device PCE for hybrid Si-PCPDTBT: PbS BHJ
solar cells with (a) different PbI2 ligand concentrations
(thickness of BHJ are all 59 nm), and (b) different BHJ film
thicknesses.

Figure 3.7: Performance of hybrid Si-PCPDTBT: PbS nanocrystal BHJ
solar cell devices with different ligands on the surface of the
PbS nanocrystals. (a) I −V curves and (b) external quantum
efficiency spectra.

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76

Figure 3.8: Time-resolved PL traces of the pristine Si-PCPDTBT polymer
and Si-PCPDTBT:nanocrystal BHJs employing NH4I- and
PbI2-exchanged PbS acceptors (λex = 500 nm; λem = 700 nm).
The instrument response function (IRF) is shown in black

78

Figure 3.9: (a) Stead-state absorption spectra of the Si-PCPDTBT thin
film before doping (black) and after a 3 min dip in a 20 ppm
solution of SbCl5 in acetonitrile (red), shown together with the
difference spectrum ( A = Adoped  Aundoped, given in blue).
(b) Transient absorption spectra of neat Si-PCPDTBT film (at

80

Figure 3.3: Normalized vis−NIR absorption spectra of PbS nanocrystals
with different surface ligands before and after 30 d aging
under ambient conditions. Top spectra are for PbI2-
exchanged PbS nanocrystals, while the bottom spectra are
for oleate-passivated PbS nanocrystals.

72

Figure 3.4: (a) ESI-MS spectrum of PbI2 dissolved in DMF solution (100
g mL
1
). The peak at 136.9 m/z indicates the presence of
I

, while the peaks at 586.7 and 588.7 m/z indicate the
presence of PbI3

. (b) Normalized vis-NIR absorption
spectra of PbS nanocrystals with different surface ligands
before and after 30 d aging under ambient conditions. Top
spectra are for PbI2-exchanged PbS nanocrystals, while the
bottom spectra are for oleate-passivated PbS nanocrystals.

Figure 3.5: (a) Schematic illustration of hybrid solar cells based on
ITO/PEDOT:PSS/Si-PCPDTBT: PbS nanocrystal/ZnO/Al
device structures. (b) Normalized absorbance spectra of PbS
nanocrystal film (black), of Si-PCPDTBT film (red), and of
a hybrid Si-PCPDTBT: PbS BHJ film (blue).

73

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xiii
0.1 ps, pumped at 660 nm) overlapping with chemically doped
steady-state spectrum of the oxidized Si-PCPDTBT film.
Steady-state absorption spectra were taken in an integrating
sphere and thus account for scattering and reflections.

Figure 3.10: Transient absorption spectra (at 4 ps, pumped at 920 nm)
overlapping with the absorption and emission spectra of the
PbS nanocrystals. (a) Hybrid Si-PCPDTBT: nanocrystal BHJ
film with PbI2-exchanged PbS, and (b) hybrid Si-
PCPDTBT:nanocrystal BHJ film with NH4I-exchanged PbS.

Figure 3.11: Transient absorption spectra of a neat Si-PCPDTBT film
(pumped at 660 nm) (a), and hybrid Si-PCPDTBT:
nanocrystal BHJ films with (b) PbI2- and (c) NH4I-exchanged
PbS acceptors (pumped at 920 nm) at different time delays
between the pump and the probe pulse.

Figure 3.12: Femtosecond TA spectra of three films as labeled. Hybrid BHJ
films were pumped at 920 nm and probed at 1200 nm; the neat
Si-PCPDTBT film was pumped at 660 nm and probed at 1200
nm.

Figure 3.13: Concentration-normalized steady-state PL spectra of PbS
nanocrystal suspensions in 1,2-dichlorobenzene with different
surface ligands: oleate (OA), NH4I, and PbI2 ( ex = 900 nm).
The feature around 1140 nm is due to an instrument artifact
from a grating change in the fluorimeter.

Figure 3.14: (a b) AFM topological image of a hybrid Si-PCPDTBT:
nanocrystal BHJ film with NH4I- and PbI2-exchanged PbS,
respectively. Image was obtained using 10  10 μm window.
(c d) TEM image of a hybrid Si-PCPDTBT: nanocrystal BHJ
film with NH4I- and PbI2-exchanged PbS nanocrystals (d = 3.2
nm). White circles in (c) indicate nanocrystal aggregates with
domain sizes ranging from 10 to 40 nm. While in (d),
individual nanocrystals can be seen embedded in the polymer
matrix.

Figure 4.1: (a) UV−vis absorption spectra of NHC−Ag, NHC−Ag2S, and
NHC−Ag2Se NC suspensions in toluene. The inset is a
photograph of dilute toluene suspensions the NCs. (b) Powder
XRD patterns of NHC−Ag, NHC−Ag2S, and NHC−Ag2Se

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xiv
NCs. (c−e) TEM micrographs of NHC-stabilized (c) Ag, (d)
Ag2S, and (e) Ag2Se NCs. The insets are the size distributions
of the corresponding NCs (n = 500 counts).

Figure 4.2: (a)
1
H NMR spectrum of NHC−AgBr, NHC−Ag NCs, and
NHC−Ag2S NCs. Resonances from 0.2 to 8 ppm are assigned
accordingly. Solvent impurities are indicated by * (CH2Cl2), Δ
(toluene), ◊ (ethanol), and ● (H2O). (b−d) High-resolution N
1s XPS spectra of NHC−Ag, NHC−Ag2S, and NHC−Ag2Se
NCs, respectively.

Figure 4.3: DOSY spectra of (a) NHC-AgBr, (b) Ag NCs, and (c) Ag2S
NCs in CDCl3 solution. Diffusion coefficients were measured
to be 6.5  10
–10
, 5.0  10
–10
, and 2.3  10
–10
m
2
/s for NHC-
AgBr, Ag NCs and Ag2S NCs, respectively.

Figure 4.4: High-resolution XPS spectra of (a) NHC-Ag, (b) NHC-Ag2S,
and (c) NHC-Ag2Se NCs, respectively. Note the absence of
Br 3d signals from NHC-Ag2S and NHC-Ag2Se NCs,
indicating the covalent interaction between the NHC ligands
and Ag2E NC surface.

Figure 4.5: (a) FT-IR and (b) TGA spectra of NHC-AgBr, NHC-Ag, NHC-
Ag2S, and NHC-Ag2Se NCs.

Figure 4.6: Geometries and binding energies for three types of binding of
an NHC ligand to (a−c) a Ag cluster that was decorated with
four Br atoms surrounding the binding site and (d−f) a Ag2S
cluster. Ag, S, Br, H, C, and N atoms are depicted in gray,
yellow, brown, pink, deep blue, and light blue, respectively.
The red Ag atom denotes the surface binding site for the NHC
ligand.

Figure 4.7: (a) Density of states for the vertex binding geometry of the
NHC molecule binding to the Ag cluster decorated with Br
atoms (upper panel) and the isolated NHC molecule (lower
panel). The dashed line denotes the Fermi level. The most
inner orbitals of the NHC molecule are aligned between the
two systems because they are only mildly affected when the
NHC molecule binds to the Ag cluster. The binding occurs
through the C atom of the ligand. The C orbitals remain nearly
intact in the valence band (rectangular) region, because

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110

112

113

xv
binding of the NHC molecule to the Ag cluster is weak. (b)
Spin-polarized density of states for the site-2 binding
geometry of the NHC molecule binding to the Ag2S cluster
(upper panel) and the isolated NHC molecule (lower panel).
The dashed line denotes the Fermi level. The C orbitals of the
ligand are significantly modified in the valence band
(rectangular) region due to strong NHC-Ag2S interaction.

Figure 4.8: (a) Powder XRD patterns of NHC-Cu, Cu2–xS, and Cu2–xSe
NCs. The asterisk in the diffraction pattern for the Cu NCs
indicates the presence of trace Cu2O. (b) UV-vis-NIR
absorption spectra of NHC-Cu, Cu2–xS, and Cu2–xSe NC
suspensions in TCE. Inset is a photograph of dilute
suspensions of Cu, CuS, CuSe, Cu1.8S and Cu1.8Se NCs in
TCE. (c-f) TEM micrographs of NHC-CuS (d = 11.7  1.8
nm), CuSe (d = 9.4  1.3 nm), Cu1.8S (d = 7.4  0.8 nm) and
Cu1.8Se (d = 9.5  1.1 nm) NCs, respectively.

Figure 4.9: Geometry and binding energy for the vertex binding of an
NHC ligand to the Cu cluster which was decorated with four
Br atoms. Cu, Br, H, C and N atoms were depicted in orange,
brown, pink, deep and light blue. The red atom denotes the Cu
atom at binding site.

Figure 5.1: Characterization of the NHC-Ag2S QDs after a 60 min
reaction: (a) powder XRD pattern; (b) UV-vis-NIR absorption
spectrum; (c) size and size distribution (300 counts for each
sample) from three different synthetic batches; (d f) TEM
micrographs. Inset in (f) is an HRTEM image showing d
(1
̅
21)
=
0.26 nm.

Figure 5.2: (a,b): UV-vis-NIR absorption spectra of in-situ reaction of
Ag2S QDs (from 1-b-AgBr) and Cu2–xS (from 1-b-CuBr),
respectively. (c e) TEM micrographs of Ag2S QDs at various
time points. Size analysis reveals d = 7.8  0.8 nm (5 min), to
8.2  0.7 nm (30 min), and 10.3  0.6 nm (60 min) (300
counts for each). (f h) TEM micrographs of Cu2–xS QDs at
different time points. Size analysis presents d = 6.8  1.0 nm
(5 min), to 8.4  1.2 nm (30 min), and 8.8  0.8 nm (60 min)
(300 counts for each).

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xvi
Figure 5.3: (a,b) TEM micrographs of Ag2S QDs prepared from Ag
(oleate), and AgBr(oleylamine) under ambient conditions,
respectively

Figure 5.4: Structures of benzimidazole- and imidazole-based NHC-AgBr
precursors.

Figure 5.5: XRD patterns of the purified products prepared from different
benzimidazole-based NHC-AgBr complexes (j-b-AgBr, j =
15) (a), and imidazole-based NHC-AgBr complexes (j-i-
AgBr, j = 1 3) (b). All of the products are phase-pure
monoclinic Ag2S (PDF#00-014-0072).

Figure 5.6: (a d) TEM micrographs of Ag2S QDs prepared from 1-b-
AgBr (10.3  0.6 nm), 2-b-AgBr (9.7  0.6 nm), 3-b-AgBr
(9.2  1.0 nm), and 4-b-AgBr (9.6  1.0 nm) complex,
respectively. The insets on the upper right corner of each
micrograph are photos of each solution mixture after a 1 h
reaction. No precipitates were observed in the reaction using
1-b-AgBr complex, while black solids were observed from
solutions using 2-b-AgBr, 3-b-AgBr, and 4-b-AgBr complex.
More aggregates were observed as the length of N-alkyl chains
decrease (from 2-b-AgBr, 3-b-AgBr, to 4-b-AgBr). The insets
on the upper left corner of (c) and (d) were representative
TEM images of corresponding Ag2S QDs, showing
aggregates upon synthesized

Figure 5.7: (a c) TEM micrographs of Ag2S QDs prepared from imidazole
-based NHC-AgBr precursors: 1-i-AgBr (4,5-H), 2-i-AgBr
(4,5-Ph), and 3-i-AgBr (4,5-Cl), respectively. The insets on
the upper left corner of each micrograph are the size and size
distributions of the corresponding QDs (300 counts each). The
inset on the upper right corner of (b) is a TEM image showing
local superstructures of Ag2S QDs prepared from 2-i-AgBr
(4,5-Ph). (d) Plot of Ag2S QD mean diameter as a function of
calculated TEP of the NHC ligands.

Figure 5.8: Characterization of surface NHC ligands on Ag2S QDs. (a)
1
H
NMR spectra of NHC-AgBr and NHC-Ag2S QDs in CD2Cl2
(from 1-b-AgBr). Resonances from 0.5 to 8 ppm are assigned
accordingly. Solvent impurities are indicated by  (toluene), *
(CH2Cl2) and  (H2O). (b,c)
1
H 
13
C HSQC spectra of NHC-

132

133

134

134

137

138

xvii
AgBr and a colloidal suspension of Ag2S QDs in CD2Cl2,
respectively. (d f) High-resolution XPS spectra of Ag 3d, S 2p,
and N 1s in Ag2S QDs, respectively.

Figure 5.9: High-resolution XPS spectra of NHC-AgBr (1-b-AgBr, a c),
NHC-Ag2S QDs (dg), and NHC-Cu2S QDs (hk). The
absence of the strong Cu
2+
satellite peaks (at 942 eV and 962
eV) in (h) proves that the oxidation state of Cu2–xS NCs is
mostly Cu
+
.

Figure 5.10: (a) FT-IR spectra and (b) TGA traces of NHC-AgBr (1-b-
AgBr), NHC-Ag2S, and NHC-Cu2–xS QDs.

Figure 5.11: Characterization of NHC-Cu1.8S QDs after a 60 min reaction:
(a) powder XRD pattern; (b) UV-vis-NIR absorption spectrum
exhibiting a broad LSPR peak; (c e) TEM micrographs of
Cu1.8S QDs. Inset in (e) is an HRTEM image showing d
(220)
=
0.20 nm.

Figure 5.12: (a) Powder XRD patterns of Ag2Se (peaks are assigned based
on calculated tetragonal Ag2Se phase) and Cu1.8Se QDs; (b)
UV-vis-NIR absorption spectra of Ag2Se and Cu1.8Se QDs;
(c,d) TEM micrographs of Ag2Se and Cu1.8Se QDs,
respectively. Insets are the HRTEM images of Ag2Se and
Cu1.8Se QDs.

Figure 6.1: Spectroscopic characterizations of Cu3 xP and InP NCs after
cation exchange: (a) Powder XRD patterns; (b) vis-NIR
absorption spectra; (c) Raman spectra. (d, e) TEM micrographs
of Cu3 xP and InP NCs, respectively. Insets are the HRTEM
images of Cu3 xP and InP NCs.

Figure 6.2: EDX spectrum of Cu3 xP (a) and InP (b) NCs, respectively.
EDX data gives a Cu:P ratio of 2.76 for Cu3 xP NCs, and a In:P
ratio of 0.76 for InP NCs. EDX analysis also reveals around
10% Br in InP NCs, which might be associated on the surface
of InP NCs.

Figure 6.3: Raman spectra of Cu3 xP and InP NCs.

Figure 6.4: High-resolution XPS spectra of Cu 2p (a), P 2p (b) in Cu3 xP
NCs and In 3d (c) and P 2p (d) in InP NCs. The absence of the

139

140

141

143

167

167

169

xviii
strong Cu
2+
satellite peaks (at 942 eV and 962 eV) in (a)
indicates that the oxidation state of Cu in Cu3–xP NCs is mostly
Cu
+
. The additional peak at 133.7 eV in (b) proves the
oxidation of P in Cu3–xP NCs.

Figure 6.5:
31
P NMR spectra of (a) a InP NCs suspension, (b) TOP, and (c)
TOPO in d6-benzene. The broad peaks at 44.4, and 26.8 ppm
can be assigned to TOPO (42 ppm) and TOP ligands ( 33
ppm), respectively. Additional sharp peaks from the InP
suspension might be from impurities in TOP/TOPO chemicals
and they are not associated on the surface of InP NCs

170

171

xix
List of Schemes

Scheme 4.1: Synthesis of NHC-stabilized colloidal coinage metal and metal
chalcogenide NCs

Scheme 4.2: General Synthetic Scheme

Scheme 5.1: Synthesis of (a) NHC-AgBr complex and (b) NHC-Ag2S QDs

Scheme 5.2: General Synthetic Scheme

Scheme 6.1: Synthesis of NHC-Cu3 xP NCs and the subsequent cation
exchange reaction with In
3+
/TOP.

104

118

129

145

165

xx
Abstract
Semiconductor nanocrystals, or quantum dots, are attractive functional materials for
photovoltaics, photocatalysis, displays, and biomedical applications because of their
uniquely tunable optoelectronic properties. The widespread implementation of these
nanomaterials strongly relies on the sophisticated manipulation of their size, morphology,
and surface functionality. Understanding nanocrystal surface chemistry and the
nanocrystal-ligand interaction is particularly important as they could affect both the
synthetic nanocrystal morphology and the optoelectronic properties. With this in mind,
my Ph.D. work has been focused on designing new inorganic/organic surface ligands,
and exploring the correlation between the surface chemistry and the charge transport
properties of quantum-dots-based photovoltaic devices, as well as understanding how the
nanocrystal-ligand interaction manipulates the resulting nanocrystal morphology.

Given the low absorption efficiency as one of the limiting factors for current state-of-the-
art hybrid polymer:nanocrystal solar cells, we sought to design the tandem and triple-
junction devices for hybrid P3HT: CdSe nanocrystal solar cells. A combination of
nanocrystalline ZnO and pH-neutral PEDOT: PSS was used as the interconnecting layer
to serve as the Ohmic contact, and the thicknesses of subcells were optimized with the
guidance of optical simulations. The tandem hybrid solar cells give a 2-fold increase in
VOC over that of single-junction devices, and light absorption in the region of 480 600
nm is increased from 60% to 80% compared to the single-junction hybrid solar cells.
Moreover, the average power conversion efficiency (PCE) was improved significantly
from 2.0% (VOC = 0.57 V) in single-junction devices to 2.7% (champion 3.1%, VOC =
1.28 V) in tandem devices and 2.3% (VOC = 1.98 V) in triple-junction devices. The results
clearly indicate that tandem/multijunction device geometries are an effective approach to
harvest more light and improve PCE in hybrid BHJ solar cells.

Another strategy to enhance absorption efficiency is to extend the semiconductor
nanocrystal absorption threshold to near-infra (NIR) region which contains about 54% of
solar energy. One of the most promising candidates is PbS quantum dots, in which the
surface chemistry is much less understood as compared Cd-based quantum dots. For
instance, it can be especially difficult to obtain colloidally stable suspensions of PbS
nanocrystals ligand exchanged with small ligands, and many atomic ligands require
dispersion in solvents that are incompatible with polymer solubility. We have developed
a hybrid iodide-based surface chemistry, which affords air- and colloidally-stable PbS
quantum dots in a polymer-compatible solvent. The modified PbS quantum dots offer a
more efficient charge separation process when blended with the donor polymer (polymer
poly[2,6-(4,4’-bis(2-ethylhexyl)dithieno[3,2-b:2’3’-d]silole)-alt-4,7-(2,1,3-
benzothiadiazole)] (Si-PCPDTBT)), as evidenced by both photovoltaic performance and
spectroscopic characterizations (i.e., transient absorption). The optimized Si-PCPDTBT:
xxi
PbS hybrid solar cells give a broad spectral response from the visible region into the NIR,
leading to a PCE of over 4% under AM 1.5G illumination.

Additionally, we have expanded the boundry of semiconductor nanocrystal surface
ligands from traditional carboxylic, phosphine, amine, thiol ligands to carbene ligands.
We designed two synthetic approaches to N-heterocyclic carbene-capped semiconductor
nanocrystals, which allows me to elucidate the energetic and dynamic nature of the
obscure carbene-nanocrystal interface across different material types. It is revealed that
these neutral L-type N-heterocyclic carbene ligands do provide strong, yet labile, binding
to various coinage-metal-based semiconductor nanocrystals. As compared to coinage
metal nanocrystals, these N-heterocyclic carbene ligands display stronger binding affinity
to semiconductor nanocrystals without introducing any sub-bandgap trap states at the
same time. This has further led to a more precise control of semiconductor nanocrystal
morphology under ambient conditions via tailoring the electronic and steric properties of
the carbene ligands. For instance, we have demonstrated that the size of Ag2S NCs
correlates linearly with the electron donating property of the NHC ligands (i.e., Tolman
electronic parameter (TEP)), with higher electron donating NHCs producing smaller
Ag2S NCs. These interesting results have also intrigued us to expand the material
platform to metal phosphide nanocrystals. We subsequently presented a facile synthetic
route towards N-heterocyclic carbene stabilized colloidal plasmonic Cu3 xP nanocrystals,
which can serve as a parent platform to access metastable wurtzite InP quantum dots that
is unusual to obtain in other traditional colloidal synthetic methods.

1
Chapter 1. Hybrid Polymer: Nanocrystal Solar Cells*
*Submitted for the book chapter in World Scientific Handbook of Organic
Optoelectronic Devices, Vol Ⅳ Organic Photovoltaics.

1.1. Introduction
Solar cells have been considered one of the most promising approaches to addressing the
global energy and environmental crisis as sunlight is an abundant, clean, and sustainable
energy source. Conventional silicon solar cells usually involve energy-intensive
fabrication processes to generate high-purity Si, resulting in high manufacturing cost. In
the search of the cost-effective materials and devices, organic photovoltaics (OPVs) have
become a very important component in third generation solar cell technologies because of
their potential to be fabricated in inexpensive ways through solution-based processing
(e.g., rapid roll-to-roll printing) with a very thin absorber layer.
1-3
Compared to traditional
inorganic semiconductors that generate free charge carriers upon photoexcitation, typical
organic semiconductors possess relatively low dielectric constants (~3 5), leading to the
formation of a Frenkel exciton. Thus, a donor-acceptor bulk heterojunction (BHJ) with a
thermodynamic driving force for exciton dissociation is required to separate the exciton
in OPVs. Although significant progress has been made in the last two decades in
developing organic donor materials (Figure 1.1), giving power conversion efficiencies
(PCEs) of single-junction OPVs up to ~10%,
4-6
the acceptor materials are still dominated
by soluble fullerene derivatives (e.g., [1]-phenyl-C61-butyric acid methyl ester
(PC60BM), [6,6]-phenyl-C71-butyric acid methyl ester (PC70BM)) since 1992.
7
This has
led to an increased emphasis on exploring new acceptor types since the optimization of
the organic donor phase may be nearing its limit to further increase device performance.
8,
2
9
For instance, one of the most important drawbacks of PCBM is the low absorption
coefficient for the visible and NIR regions of the solar spectrum. As a direct
consequence, significant transmissive losses occur in single-junction OPVs because the
organic donor material itself only displays a relatively narrow absorption band.
Furthermore, PCBM often suffers from chemical and photo-induced degradation under
ambient conditions, impeding the potential utilization for real-world applications.
Therefore, developing electron acceptor materials with higher absorption coefficients
over a wider spectral range, and higher optical and chemical stabilities is very much
desirable.
Figure 1.1 Chemical structures of conjugated polymers discussed in this chapter.

In the place of fullerene derivatives, colloidal inorganic nanocrystals (NCs) have
arrived as a promising alternate electron acceptor,
10
with the aim of combining the
beneficial properties of both organic and inorganic phases. Compared to the organic
phase, inorganic semiconductor NCs possess intrinsically higher electron mobilities,
broad and tunable absorption with relatively high absorption coefficients, better
3
stabilities, and higher dielectric constants. These attributes are potentially favorable with
respect to several mechanistic events in solar cell operation. First, the high dielectric
constant ( r > 10) of inorganic NCs reduces Coulombic interaction, resulting in
delocalized excitons (Wannier-Mott excitons), which in turn facilitates charge separation
in the donor-acceptor charge transfer (CT) interface. Additionally, the dimensionality
tuning of semiconductor NCs can be used to improve solar cell device performance
through morphological control, with anisotropic nanorods (NRs) and branched multipods,
such as tetrapods (TPs), aiding directional electron transport along their principal rod axis
and reducing the inefficient electron hopping between NCs. Moreover, inorganic
semiconductor NCs can be readily synthesized with different chemical composition (e.g.
Ⅱ Ⅵ vs Ⅳ Ⅵ), surface ligand, and size, leaving a vast parameter space for device
optimization from the perspective of acceptor materials. These materials have led to the
growing field of so-called hybrid polymer: nanocrystal solar cells.
In this chapter, we will provide the most up-to-date overview of the progress
made in polymer: nanocrystal hybrid solar cells, along with discussion of the most
important mechanistic processes in hybrid devices, which will give insight into further
improving device efficiencies. We focus the review mainly on solar cells containing both
solution-processable polymer and colloidal NCs as the photoactive layer of devices,
excluding dye/quantum dot sensitized solar cells, colloidal quantum dots solar cells, and
hybrid perovskites, for which processes are fundamentally different and extensive
literature reviews can be found elsewhere. Here, we first introduce semiconducting NCs
and their advantageous properties for photovoltaic applications. Then, we discuss the
mechanism and theory of hybrid solar cells, highlighting the limiting factors for optimal
4
device performance. Subsequently, we present various approaches to improve device
performance of hybrid solar cells, focusing on engineering surface ligands, tuning NC
shapes and film morphology, and optimizing device architectures. A new direction
involving inorganic NIR absorbers, such as the lead chalcogenide NCs, is then reviewed,
followed by an outlook of polymer: nanocrystal hybrid solar cells.

1.2. Introducing Inorganic NCs
Inorganic semiconductor NCs, also called quantum dots (QDs), are highly attractive
materials for photovoltaics because of their uniquely tunable optoelectronic properties,
rising from quantum confinement effects. Quantum confinement can be observed when
the particle size is of the same magnitude as the de Broglie wavelength of the electron
wave function, in which discontinuous, atomic-like electronic states are formed. The
band gap becomes size dependent, and it increases with decreasing particle size.
Therefore, the absorption spectrum of semiconductor NCs can be readily tailored to
various spectroscopic windows by simply tuning the particle size and shape without
changing the chemical composition.
11

Additionally, the development of robust synthetic methodologies has led to
semiconductor NCs with well-defined morphology (i.e., size, size distribution, shapes)
and surface chemistry. Modern spectroscopic techniques show that colloidal NCs are
usually decorated with organic molecules (ligands) on their surface, which play an
essential role in processing and optoelectronic device performance.
12, 13
These long-chain,
insulating organic ligands are inherent to NCs synthetically, and give colloidal stability,
as well as surface passivation of trap states, or dangling bonds. The nature of surface
5
ligands can affect trap states and trap state density, frontier energy levels (through
quantum-confined Stark effect), and interfacial electronic coupling,
14, 15
all of which are
closely related to exciton separation and charge transport in photovoltaic devices. For this
reason, colloidal semiconductor NCs should be considered as a NC/ligand hybrid system,
wherein the interface needs be carefully addressed for photovoltaic applications.
Introducing semiconductor NCs in photovoltaic devices should, in theory, be able
to surpass the inherent Shockely-Queisser limit (34% for the optimal 1.34 eV band gap)
for single-junction solar cells.
16
There are two main mechanistic losses based on the
Shockely-Queisser calculation: (1) spectral loss (h < Eg), and (2) hot carrier loss (i.e.,
thermal loss, h > Eg), which sums to 52% energy loss in total.
17
Spectral loss can be
minimized by building tandem junctions using materials with complementary absorption
spectra, while thermal loss could be overcome by the so-called multiple exciton
generation (MEG). When electrons and holes are excited with photon energies larger than
the band gap energy, they become hot carriers and often relax rapidly to band edges via
phonon emission. This Auger-mediated process was found to be slowed in quantum-
confined materials, such as semiconductor NCs, where relatively sparse electronic states
are available.
18
When the hot excitons bear more than twice the energy of the band gaps
of semiconductor NCs (h > 2 Eg), they can relax to their band edges by exciting an
adjacent electron from the valence to conduction band, resulting in two excitons per
photon. Thereby, this process provides the possibility to minimize thermal loss and
consequently overcome the Shockley-Queisser limit.
19
Although the concept of MEG in
photovoltaics has brought intensive discussion since 2004,
20
experimental proof of MEG
in working devices remain mostly spectroscopic.
21-24
It is recently revealed that the
6
enhancement of MEG process in semiconductor NCs-based photovoltaics strongly
depends on the chemical composition (e.g., heterostructures), morphology, and surface
chemistry of the semiconductor NCs.
25, 26
The concept of utilizing semiconductor NCs to
reduce thermal loss through MEG is still very promising.

1.3. Hybrid Polymer: Nanocrystal Solar Cells
1.3.1. Mechanism and Theory
Much like the conventional polymer-fullerene OPVs, hybrid polymer: nanocrystal solar
cells require BHJs to facilitate the dissociation of singlet excitons from photo-excited
polymers due to the strong binging energy of these excitons (Eb = 0.1 0.4 eV). In hybrid
BHJs, where conjugated polymer and semiconductor NCs percolate throughout the
photo-active layer, the polymer phase serves as an electron donor, while the
semiconductor NCs work as electron accepters (Figure 1.2 a). When photons are
absorbed by the hybrid active layer, excitons are generated in the conjugated polymers,
which diffuse to the polymer-NCs interface, where they can then dissociate into free
charges, driven by the type-Ⅱ energy alignment between polymer and NC. Subsequently,
electrons are transported through the NC percolation network to the cathode, while holes
are transported via polymer pathways to the anode. Meanwhile, photo-excited
semiconductor NCs can also generate excitons, which are dissociated by hole transfer
from the NCs to the polymer, additionally contributing to photocurrent in hybrid devices.
Based on the empirically determined band gaps and exciton binding energies for
conventional donor polymers and CdSe acceptor NCs, as well as very conservative
estimates for the external quantum efficiency EQE (i.e., 40% EQE for polymer and 30%
7
EQE for CdSe NCs) and fill factor (FF = 0.50), hybrid polymer: CdSe BHJ devices are
believed to be able to achieve PCEs over 10% (Figure 1.2 b).
27, 28
This is not the
theoretical efficiency limit for hybrid BHJs, but rather a fair estimation of the
performance that these devices should be able to achieve using current materials
platforms. However, the highest reported PCEs for hybrid BHJs to date remain in the
range of 5 6%,
29
lagging behind those values. A better understanding of the origin of this
photovoltaic performance gap from a mechanistic standpoint is thus needed.

Figure 1.2 (a) Schematic illustration of hybrid polymer: NC BHJ solar cells and type-Ⅱ energy
level diagram, showing the dissociation of excitons and charge transport process. (b) Calculated
power conversion efficiencies of a hybrid polymer: CdSe solar cell as a function of the optical
band gap of the donor material and CdSe NC diameter. Adapted with permission from [28]. 
2015 Elsevier.

In a typical excitonic solar cell, the complete physical process comprises six
fundamental steps: (i) photon absorption ( a), (ii) exciton generation ( ex), (iii) exciton
diffusion ( diff), (iv) exciton dissociation ( ed), (v) charge transport ( tr), and (vi) charge
collection ( coll), where  is the associated yield of each process. The EQE is thus
determined by the product of each yield, EQE = a  ex  diff  ed  tr  coll. To
8
better probe these mechanistic processes in hybrid devices, we sought to estimate the
efficiency of these processes experimentally for the hybrid PCPDTBT: CdSe NCs BHJs
(PCPDTBT = poly[2,6-(4,4′-bis(2-ethylhexyl)dithieno[3,2-b:2′3′-d]silole)-alt-4,7-(2,1,3-
benzothiadiazole)]).
30, 31
We first simplified them into four mechanistic steps, (i) photon
absorption ( a), (ii) exciton dissociation ( ed), (iii) charge separation ( sep), and (iv)
charge collection ( coll). By using ultrafast transient absorption spectroscopy, we were
able to detect unambiguous spectral signatures of interfacial charge transfer (ICT) states.
Global analysis of these signatures over different time periods further allowed us to
quantify the yield of these processes, which are discussed in detail in the following,
elucidating the major bottlenecks for hybrid BHJ solar cells.
The photon absorption efficiency a depends on the absorption spectral band,
optical absorption coefficients, thickness of the photoactive materials, and the device
configuration. The development of new low band gap conjugated polymers (e.g.,
(PCPDTBT) and 2,6-(N-(1-octylnonyl)dithieno[3,2-b:20,30-d]pyrrole)-alt-4,7-(2,1,3-
benzothiadiazole)) (PDTPBT), Figure 1.1) with higher absorption coefficients as donor
materials can significantly stimulate the photovoltaic performance of hybrid BHJs.
Similarly, the utilization of NIR-absorbing semiconductor NCs, such as PbS and PbSe, as
electron acceptors could also boost the photocurrent due to the considerable contribution
from the NIR response. Furthermore, by utilizing the optimal device configuration, such
as conventional or inverted device architectures, or by building tandem junctions, a can
also be improved due to the enhanced optical field intensity. It is noteworthy that photon
absorption efficiency is calculated to be one of the limiting factors of these hybrid BHJs,
with an a of only ca. 17.5%
30
in the hybrid PCPDTBT: CdSe NCs BHJs.
9
Dissociation of the exciton into free carriers at the polymer-NC interface is the
most interesting step for hybrid solar cells. The exciton dissociation efficiency ( ed) is
correlated to the donor-acceptor (D/A) LUMO energy offsets, which can be described by
Marcus Theory.
30, 32
It is generally believed that the minimum energy offset required to
dissociate an exciton at the D/A interface is the energy to overcome the exciton binding
energy of the polymers (0.1 0.4 eV). Apart from the choice of polymer and
semiconductor NC, the surface ligands on the NCs also play a crucial role in ed. First,
there is a dramatic effect on the electronic coupling between the polymer and NCs
through spatial coupling, resulting from the inherently large surface area to volume ratio
of the NCs. It was first demonstrated that efficient charge transfer occurs in poly[2-
methoxy-5-(2-ethylhexyloxy)-1,4-phenylenevinylene] (MEH-PPV) blended with CdS
and CdSe NCs in 1996.
34
In the reported hybrids, efficient charge transfer only occurred
in MEH-PPV/pyridine-exchanged NCs, and no charge transfer was observed in MEH-
PPV/TOPO-capped NCs, showing the importance of surface ligands on the exciton
dissociation efficiency. By selectively exciting the donor polymer phase using transient
absorption spectroscopy, we determined that electron transfer from the polymer
PCPDTBT to tert-butylthiol-capped CdSe NCs happens on an ultrafast timescale (< 65
fs), with an estimated electronic coupling between the polymer phase and NCs to be J 
17 meV (Figure 1.3a).
30
This timescale for electron transfer is comparable to polymer:
fullerene BHJs.
35
The amplitude of the spectral bleach from the reduced QD acceptors
allows for direct assessment of the exciton dissociation yield, ed, to be 82  5%.
Therefore, the electron transfer (exciton dissociation) process occurs rapidly with
relatively high yield, and it does not appear to be the major limitation in hybrid solar
10
cells.
Figure 1.3 (a) Transient absorption spectra of a neat CdSe QDs film (yellow), a neat PCPDTBT
film (blue) and a hybrid PCPDTBT: CdSe BHJ film (green) (  ex = 800 nm). The negative spectra
signature at 610 nm indicates that the quantum dot 1S e energy level is populated by charge
transfer from PCPDTBT. (b) Density of electrons located on the NC conduction band as a
function of time, attained from the deconvolution of the 610 nm transient absorption bleach.
Adapted with permission from [30].  2013 American Chemical Society.

Due to the low dielectric constants of conjugated polymers, dissociated excitons
need to overcome their mutual Coulombic attraction in the ICT states and separate into
free carriers, which can then be transported to the electrodes. Based on the calculation of
electron density located on the NC conduction band as a function of time, we determined
the charge separation efficiency to be sep = 29.5  4.5% (Figure 1.3 b), suggesting that
the rapid and measurable geminate recombination between NC electrons and polymer
polarons is a major limitation of hybrid BHJs. This has been further confirmed by
spectroscopic evidence
36-38
and computational simulations.
39
It is generally believed that
the NC surface traps strongly contribute to the electron hole recombination.
40
By
comparing the charge carrier dynamics between P3HT: PCBM and hybrid P3HT: CdSe
BHJs using transient photovoltage and photocurrent measurements, McNeil et al.
concluded that surface traps on NCs lead to a low yield of charge extraction of the long-
lived carriers.
40
This is further confirmed by Greenham et al., showing that trap-mediated
11
recombination is an essential factor in the poor charge separation efficiency in hybrid
P3HT: CdSe BHJs.
41
However, computational simulation by time-domain ab initio
analysis reveals that the electron hole recombination occurs several orders of magnitude
faster in the polymer than in the NC phase due to the much higher, bulk-like density of
electronic states in NCs.
39
While the polymer is continuous in space, the electronic states
are molecule-like and discrete in energy. Therefore, the charge transfer to NCs is much
faster than to the polymer, and the electron hole recombination occurs faster in the
polymer phase. For this reason, the authors attributed the significant electron hole
recombination to the highly inhomogeneous energy landscape of electronic states in the
polymer NC interface. Although additional investigations into the recombination
mechanism would still be highly valuable, the current studies suggest that the density of
surface trap states of semiconductor nanocrystals needs to be optimized through ligand
engineering to minimize the charge recombination at the ICT states.
Once separated, the charges are then transported to electrodes as free carriers. The
carrier transport efficiency is based on the carrier mobility of donor and acceptor phases,
the film morphology, and the surface trap states. For instance, efficient electron transport
between semiconductor NCs requires that the NCs be physically touching each other,
thus forming a percolation pathway with a relatively low density of trap states. The
employment of anisotropic semiconductor NCs, such as rods and tetrapods, has been
shown to significantly enhance the electron transport efficiency in the hybrid BHJs due to
a reduced number of inefficient “electron-hopping” events between NCs.
10,42

Furthermore, the inefficient “electron-hopping” events between NCs can be improved by
surface ligands via improved interparticle spatial coupling. For example, after ligand
12
exchange with tert-butylthiol for CdSe NCs, we calculated the lower limit for charge
collection efficiency, coll, to be 66  14% in the hybrid PCPDTBT: tert-butylthiol-
exchanged CdSe solar cells. Most recently, Talapin et al. demonstrated that a thin film
made of atomic-halide-capped CdSe NCs possesses an electron mobility as high as e =
12 cm
2
V
–1
s
–1
, which is several orders of magnitude higher than that of an oleate-capped
CdSe NC film.
43
By passivating the surface traps of CdSe NCs with indium, Kagan et al.
achieved a high electron mobility of e = 18 cm
2
V
–1
s
–1
from CdSe films.
44
As for the
hybrid polymer: nanocrystal solar cells, Fu et al. demonstrated that by tuning the
benzenedithiol ligand orientation, hybrid PCPDTBT: CdSe NC solar cells were able to
achieve a high PCE of 4.18%, owing to the significantly improved electron mobility of
CdSe NCs.
45

Altogether, current spectroscopic analyses indicate that the main reasons for poor
device performance in hybrid BHJs are low absorption yield and limited charge
separation efficiency resulting from trap-mediated electron hole recombination. Here, we
summarize recent progress in addressing the aforementioned limitations, highlighting the
importance of surface ligands present on NC surafaces with respect to the charge transfer
dynamics, and discuss the crucial influence on both the short-circuit current (JSC) and the
open circuit potential (VOC) in hybrid BHJ solar cells.

1.3.2. Surface Chemistry of Colloidal Semiconductor NCs
In order to manipulate surface traps and electronic coupling between the polymer-NC and
NC-NC interfaces in hybrid BHJ solar cells, the surface chemistry of the acceptor NCs
must be thoroughly understood and rationalized. Monodisperse colloidal semiconductor
13
NCs are generally synthesized with the help of insulating, long-chain aliphatic organic
ligands, which can control reaction kinetics and particle growth. These organic ligands
further give post-synthetic colloidal dispersibility in nonpolar solvents (such as toluene,
hexanes) through van der Waals stabilization. Therefore, as-synthesized semiconductor
NCs inevitably contain long-chain insulating organic ligands, such as amines, carboxylic
acids, phosphine oxides, thiols, and phosphonic acids. These ligands are traditionally
categorized as neutral L-type ligands (Lewis bases, dative covalent bonding, e.g., amine,
phosphine, or thiol), anionic X-type ligands (normal covalent bonding, e.g., carboxylate,
phosphonate, or thiolate), and Z-type ligands (Lewis acids, e.g., Cd(O2CR)2)
46
(Figure
1.4).
Figure 1.4 Schematic illustration of surface ligand coordination chemistry according to the
covalent bond classification method. Adapted with permission from [46].  2013 American
Chemical Society.

The nonstoichiometric nature at the surface of many semiconductor NCs is
exacerbated in colloidal NCs due to the large surface-to-volume ratio, causing different
types of ligands to interact with the surface atoms quite differently. In general, the as-
synthesized colloidal NCs display a metal-rich surface. It has been shown that, depending
14
on the synthetic approaches and purification procedures, CdSe NCs can display a Cd/Se
ratio as high as ~6: 1.
47, 48
The excess metal cations at the surface of the NCs thus require
anionic ligands for charge balance. That is, the surface of nonstoichiometric NCs must be
coordinated to an appropriate amount of X-type ligands, with L-type and Z-type ligands
filling in the rest of the coordination sphere of the NCs.
Ligand binding to NCs can either introduce or eliminate surface trap states. The
origin of trap states in semiconductor NCs can be well illustrated by molecular orbital
diagram.
49
As shown in Figure 1.5, if the excess surface metal cations are not coordinated
by ligands, the resulting uncoordinated surface sites will give electronic states within the
band gap via dangling bonds, acting as trap states. Similar trap states can also be
introduced if the surface chalcogens (S, or Se) are not properly passivated by strong
Lewis acids (Z-type ligands). On the other hand, if the organic ligands are weakly
coordinated to the excess metal cations on the nanocrystal surface, the resulting
bonding/antibonding states could lay within the band gap, also serving as trap states.
Ideally, the passivating ligands on the NC surface should coordinate strongly enough
such that the resulting bonding and antibonding states fall within the valence and
conduction bands, respectively, eliminating electronic states within the band gap.
Therefore, when developing new surface chemistry of NCs for hybrid photovoltaics, one
must consider the following important factors: (i) the category of ligands (L-, X- or Z-
type), (ii) the bonding nature of ligands (whether it introduces or eliminates midgap trap
states), (iii) and the landscape of the surface of the NCs (whether it is reconstructed upon
exchange).
15
Figure 1.5 Qualitative molecular orbital diagrams illustrating the bonding interactions between
surface Cd
2+
cations on CdSe NCs and σ-donating ligands. (a) An undercoordinated Cd
2+
surface
site yields a ‘nonbonding’ state within the HOMO LUMO gap, which is typically referred to as a
‘dangling bond’. (b) A strong σ-donating ligand binds to surface Cd
2+
, leading to bonding and
antibonding states laying outside of the band gap, thus eliminating surface traps. (c) A weak σ-
donating ligand forms an antibonding state within the HOMO LUMO gap, and creates a surface
trap state. Adapted with permission from [49].  2010 American Chemical Society.

1.3.3. Tailoring the Interfacial Coupling and Energetics Through Ligand
Engineering
It is worth noting that all of the high-performing, state-of-the-art hybrid BHJ solar
cells mandate the removal or displacement of the electrically insulating organic ligands.
It is generally achieved by a post-synthetic ligand exchange protocol to replace the
insulating anionic X-type ligands with smaller organic ligands. Such ligand exchange
processes are extremely critical for hybrid devices (vide supra) since they can impart
better polymer-NC and NC-NC electronic coupling, and thereby facilitate efficient
charge separation and charge transport processes.
36, 50-52
In general, the ligand exchange
can be performed either in the colloidal-solution phase or in the solid phase after the
deposition of BHJ layer (post-deposition, thin film ligand treatment). Colloidal ligand
exchange is typically conducted by adding an excess amount of the desired exchanging
ligands to a suspension of native-ligand-capped NCs, creating a concentration gradient
that drives the equilibrium of the surface moiety towards the desired ligands, assuming
proper charge balance. An appropriate antisolvent is subsequently added into the
dispersion to assist with the exchange process. This may be repeated several times to
16
optimize the extent of ligand exchange, which can be directly assessed by NMR, FT-IR,
and thermogravimetric analysis (TGA). Alternatively, solid-phase ligand exchange is
generally performed by dipping the solid BHJ film into a dilute solution containing the
replacement ligand. Common bidentate ligands (e.g., 1,2-ethanedithiol (EDT), 1,3-
benzenedithiol (BDT), 3-mercaptoproprionic acid) and atomic halide ligands (e.g.,
iodide) have been successfully applied to hybrid BHJs via solid-phase ligand exchange.
29,
53, 54
However, solid-phase ligand exchange suffers from several inherent drawbacks for
hybrid BHJs: (i) ligand exchange can be impeded by slow and/or incomplete solid-state
ligand diffusion kinetics through the BHJ film. The effectiveness of this process is also
difficult to assess experimentally, (ii) Reduction of the film volume through exchange of
small ligands for larger ones can lead to severe film cracking, and (iii) The dip-washing
process may be incompatible with large-scale solution processing. Therefore, a
quantitative colloidal ligand exchange is conceptually preferable for BHJ films; however,
it has been demonstrated that both approaches give significantly improved device
performance (vide infra). In order to tailor the polymer-NC and NC-NC interface via
ligand engineering in hybrid solar cells, it is very important to thoroughly understand the
ligand exchange chemistry. For instance, the efficacy of ligand exchange is essential,
because insulating native ligands left behind on the NC surface can negatively affect
exciton dissociation and charge transport efficiency. Moreover, the introduction or
elimination of surface trap states associated with ligand exchange procedures must be
considered, as they can potentially influence both the JSC and VOC via trap-mediated
charge recombination. Additionally, it is a prerequisite that the solvents that the
colloidally exchanged-NCs are suspended in be compatible with the polymer solubility in
17
order to fabricate hybrid BHJs.
In the context of ligand engineering in hybrid devices, extensive research has
been conducted in Cd-based hybrid BHJs because of the well-known surface chemistry of
CdE (E = S, Se) NCs. The most common ligands applied for solid phase ligand exchange
in hybrid BHJs are bidentate ligands (Table 1.1). These bifunctional ligands can decrease
the NC and polymer interfacial distance, yielding a stronger spatial and electronic
interaction between the NCs and polymer. Furthermore, they can crosslink adjacent NCs
and increase the degree of wavefunction overlap between NCs, thus leading to increased
interparticle electronic coupling. In 2011, Ren et al. achieved a record PCE of 4.1% by
grafting CdS NCs onto crystalline P3HT nanowires and exchanging surface ligands by
EDT.
55
They attributed the improved performance to the better electronic coupling
between donor-acceptor and acceptor-acceptor components. More recently, Zhou et al.
proved that EDT treated PCPDTBT: CdSe BHJs give a PCE of 4.7%.
53
In addition to the
enhanced electron mobility of NCs, they also ascribed the improved performance to the
reduction of defect state density on the NC surface. Liu et al. further applied 1,3-BDT to
PDTPBT: PbS0.6Se0.4 hybrid BHJs, and by optimizing phase segregation, they were able
to achieve the best performing hybrid solar cells with a PCE of 5.5%.
29

With regards to colloidal ligand exchange, one of the most prevalent colloidal
ligand exchange protocols for CdSe NCs is using pyridine as the replacement ligand.
56

The resulting hybrid device efficiency has been demonstrated to be dramatically
improved compared to the corresponding polymer blended with native-ligand-capped
CdSe NCs. In fact, when PCPDTBT was blended with a mixture of pyridine-capped
CdSe NCs and nanorods, hybrid solar cells with PCEs over 3% were obtained.
57, 58

18
However, the pyridine exchange protocol is generally not complete since pyridine is a
neutral and weak L-type ligand. When replacing anionic X-type native ligands (e.g.,
carboxylate ligands) that are required to balance surface charges, the pyridine exchange
process is thermodynamically unfavorable without surface restructuring through loss of
Z-type ligands CdX2,
46
leaving some deleterious insulating ligands. The improved
photovoltaic performance may be attributed to the fact that the weakly binding pyridine
generates ‘hot spots’, in which it becomes spatially possible for the charge transfer
between the donor and acceptor, even if the pyridine ligand exchange is not complete.
To this end, we sought to explore strongly binding ligands with low organic
content that can quantitatively remove the insulating ligands and result in good polymer-
NC and NC-NC coupling in hybrid devices. We showed CdSe NCs ligand exchange with
tert-butylthiol, allowing the quantitative removal of native oleate ligands, as evidenced
by TGA and FT-IR spectroscopy (Figure 1.6 a, b), while still maintaining excellent
colloidal stability and solution processability.
59
The quantitative ligand exchange with
tert-butylthiol was possible due to the fact that tert-butylthiol can coordinate to surface
Cd
2+
both as X-type (tert-butylthiolate) and L-type (tert-butylthiol) ligands. Because of
the improved NC interparticle coupling, electrochemical photocurrent measurements
displayed a >70-fold enhancement in photocurrent for tert-butylthiol-exchanged CdSe
NC films as compared to pyridine-exchanged CdSe NC films. It was further observed by
field effect transistor (FET) measurements that carrier mobility in tert-butylthiol-
exchanged CdSe NC film was substantially increased. Subsequent photoluminescence
lifetime studies suggest that the tert-butylthiol-exchanged CdSe NCs also possess a lower
density of surface trap states as compared to pyridine-exchanged CdSe NCs (Figure 1.6
19
c). This is further confirmed by photo-induced current transient spectroscopy, showing
that tert-butylthiol ligands passivate the deepest trap levels in P3HT: CdSe hybrids.
52
A
higher EQE was observed in P3HT: CdSe BHJ solar cells when tert-butylthiol-exchanged
CdSe QDs were blended as compared to pyridine-exchanged CdSe QDs. Together, these
data indicate that tert-butylthiol is a superior ligand for NC acceptors with respect to
charge transfer and carrier transport in hybrid solar cells.
Figure 1.6 Thermogravimetric analysis (TGA) traces (a), and (C H) stretching region in the
FT-IR spectra (b) of CdSe NCs with their native ligand (NL), after colloidal ligand exchange with
pyridine (Py), and tert-butylthiol (tBT). The low organic content and the absence of the high
temperature mass loss observed from TGA traces suggest a quantitative ligand exchange by tBT
ligands, which is in accordance with the relative intensity of (C H) stretching based on FT-IR
spectra. (c) PL lifetime decay traces for films of neat CdSe QDs (  ex = 400 nm,  em = 650 nm).
Lifetimes of 1.4, 3.3, and 4.4 ns were measured for CdSe NCs with NL ligands, and after
exchanged with Py and tBT ligands, respectively. Adapted with permission from [59, 60]. 
2012 American Chemical Society.

In addition to the ligand-induced enhanced electronic coupling at polymer-NC
and NC-NC interfaces, which yields increased photocurrent (JSC) in hybrid devices, we
proved that the surface ligands are also capable of tailoring the frontier valence and
conduction band edges of NCs (or HOMO/LUMO energies). Consequently, the energy
offset ( EDA, Figure 1.2 a) between the HOMO of the donor polymer and the conduction
edge (LUMO) of the NC acceptor can be systematically tuned. As the VOC of hybrid
devices is proportionally related to EDA,
61-63
tuning the LUMO energy of the acceptor
material will directly affect the VOC, which in turn, affects device performance. We
20
showed that the introduction of strongly binding and electron-donating ligands with a
small organic content, such as tert-butylthiol enables an increase of VOC in hybrid devices
by elevating the LUMO energy of the NCs and thereby increases the energy offset from
the polymer HOMO. Hybrid solar cells based on P3HT and tert-butylthiol-capped CdSe
NCs achieved PCEs of 2.0%, with a substantially high VOC of 0.80 V.
60
In comparison,
hybrid devices made from pyridine-exchanged and non-ligand exchanged CdSe exhibited
lower VOC of 0.57 V and 0.70 V, respectively. Such consistently high VOC in the thiol-
exchanged CdSe NCs is attributed to the elevated LUMO energy of NCs relative to the
HOMO energy of P3HT, which creates the largest EDA among these ligands. This is
unambiguously evidenced by cyclic voltammetry, differential pulse voltammetry, and
spectroelectrochemistry (Figure 1.7 a, b). The thiol-exchanged CdSe NCs indeed hold the
highest lying LUMO among these three ligands. It is noteworthy that tailoring the LUMO
energy of the NCs by changing the surface ligands has been reported previously,
however, we demonstrated for the first time that there is a direct correlation between the
ligand-induced LUMO energy shift and the resulting VOC in the working solar cells.
Furthermore, it appears that the increase of the energy of the LUMO of acceptor NCs
does not lead to a decrease of electron transfer rate from polymer to NCs, as
demonstrated by the dramatic photoluminescence quenching. This might be ascribed to
the high density of states in the conduction band of the acceptor NCs. Therefore, by
carefully engineering the surface ligands of NCs, one can simultaneously increase both
JSC and VOC. This tailoring of the heterogeneous interface coupling and LUMO energy of
NCs provides a principal rationale to optimize hybrid solar cells.

21
Figure 1.7 Visible absorption spectra for (a) Py- and (b) tBT-exchanged CdSe NCs on ITO
electrodes and collected with decreasing applied bias. All potentials reported here are referenced
to NHE. Initial bleaching is first observed at 0.54 and 0.74 V for Py- and tBT-exchanged CdSe
NCs, respectively. (c) Cyclic voltammograms for CdSe NCs exchanged with Py and tBT,
indicating that a higher lying LUMO for the tBT-exchanged CdSe NCs. Adapted with permission
from [60].  2012 American Chemical Society.

Table 1.1 A summary of high performing hybrid polymer: nanocrystal solar cells
Polymer NCs Morphology Ligand
J SC
(mA cm
2
)
V OC
(V)
FF
PCE
(%)
Ref.
PCPDTBT CdSe QDs & NRs HA 8.7 0.63 0.56 3.1 [64]
PCDTBT CdSe MPs tBT 10.0 0.82 0.39 3.2 [65]
P3HT CdS QDs EDT 10.9 1.1 0.35 4.1 [55]
PCPDTBT CdSe MPs tBT 11.6 0.71 0.49 4.1 [65]
PCPDTBT CdSe NRs EDT 12.8 0.74 0.50 4.7 [53]
PSBTBT-
NH 2
CdTe TPs 1,3-BDT 7.2 0.79 0.56 3.2 [66]
PDTPBT PbS QDs EDT 13.6 0.57 0.51 3.8 [67]
Si-
PCPDTBT
PbS QDs EDT 10.8 0.63 0.51 3.5 [68]
PDTPBT PbS xSe 1-x QDs 1,3-BDT 14.7 0.57 0.66 5.5 [29]
Si-
PCPDTBT
PbS QDs PbI 2 18.2 0.48 0.55 4.8 [54]
PCPDTBT = poly[2,6-(4,4-bis(2-ethylhexyl)4H-cyclopenta[2,1-b:3,4- b’]-dithiophene)-alt-4,7-(2,1,3-
benzothiadiazole)].
PCDTBT = poly[N-9’-heptadecanyl-2,7-carbazole-alt-5,5-(4’,7’-di-2-thienyl-2’,1’,3’-benzothiadiazole)],
P3HT = poly(3-hexylthiophene).
PSBTBT-NH 2 = poly[(4,4’-bis(2-ethylhexyl)-dithieno[3,2-b:2’,3’-d]silole)-2,6-diyl-alt-(2,1,3-
benzothiadiazole)-4,7-diyl].
PDTPBT = 2,6-(N-(1-octylnonyl)dithieno[3,2-b:20,30-d]pyrrole)-alt-4,7-(2,1,3-benzothiadiazole)).
Si-PCPDTBT = poly[2,6-(4,4′-bis(2-ethylhexyl)dithieno[3,2-b:2′3′-d]silole)-alt-4,7-(2,1,3-
benzothiadiazole)].
NRs = nanorods.
MPs = multipods.
HA = hexanoic acid.
TPs = tetrapods.
BDT = benzenedithiol.
EDT = 1,2-ethanedithiol.

1.3.4. Charge Transport Dependence on NC Shape and Film Morphology
As mentioned earlier, the exciton dissociation process at hybrid interfaces is in fact fast
22
and efficient, and is therefore is not the main limiting process of hybrid devices.
However, efficient charge transport to electrodes before recombination is still extremely
important. Thus, it is very critical to optimize hybrid materials and film morphology so
that percolated pathways for dissociated carriers are available.
It is worth noting that appreciable PCEs of hybrid polymer: nanocrystal solar cells
were not obtained until the use of CdSe NRs (PCE = 1.7%).
10
Subsequently, they
incorporated three-dimensional hyperbranched CdSe NCs in hybrid devices, yielding an
improved PCE of 2.2%.
69
It is generally believed that anisotropic nanostructures can
provide directional pathways for electron transport within the NCs, reducing inefficient
electron hopping events between the NCs. This concept has been applied to various
systems (Table 1.1), consistently giving superior photovoltaic performance. For example,
one of the best hybrid devices to date is using CdSe NRs blended with a low band gap
polymer PCPDTBT. After replacing the original ligands with EDT, Zhou et al. were able
to obtain an average PCE of 4.7%.
53
To further extend this concept, we demonstrated that
anisotropic tert-butylthiol-exchanged CdSe multipods could also lead to improved device
performances when blended with a series of donor polymers, including a novel semi-
random poly(3-hexylthiophene-thiophene-diketopyrrolopyrrole) (P3HTT-DPP)
copolymer, an alternating copolymer with a low-lying HOMO level poly[N-9’-
heptadecanyl-2,7-carbazole-alt-5,5-(4’,7’-di-2-thienyl-2’,1’,3’-benzothiadiazole)]
(PCDTBT), along with the more commonly used P3HT and PCPDTBT (Figure 1.1).
65

Enhanced current densities (9.5 11.6 mA cm
–2
) were obtained in hybrid devices owing to
the anisotropic nanostructure of the acceptor phase. As a result, champion PCEs as high
as 3.1, 3.2, and 4.1% were achieved in P3HTT-DPP, PCDTBT, and PCPDTBT: CdSe
23
hybrid devices, respectively. Meanwhile, PCDTBT: CdSe hybrid devices gave the
highest VOC (0.82 V) among all hybrid solar cells.
Figure 1.8 (a) TEM micrograph of CdSe multipods. (b) I V curves under dark (dashed lines) and
AM 1.5G 1 sun illumination (solid lines). (c) Ideal morphology for polymer: nanocrystal BHJs
bearing some degree of vertical phase segregation, with donor materials aggregating near the
anode and acceptor materials assembling near the cathode. (a) and (b) are adapted with
permission from [65].  2013 Royal Society of Chemistry.

Similarly, optimizing the polymer donor phase is also beneficial to hybrid device
performance. This is because most conjugated polymers used in OPVs are semi-
crystalline wire-like phases, and the crystallinity of polymers can be altered through
thermal or solvent vapor annealing. For instance, applying thermal annealing to OPVs
generally leads to improved charge transport (and thus photovoltaic performance) due to
increased hole mobility.
70-72
This highlights the importance of hole carrier mobility in the
polymer donor phase in hybrid devices. More importantly, post-deposition thermal or
solvent vapor annealing could further manipulate the nanoscale morphology of hybrid
BHJs. Ideally, the optimized nanoscale morphology for hybrid BHJs would possess a
degree of vertical phase segregation, with donor materials aggregating near the anode and
acceptor materials assembling near the cathode (Figure 1.8 c). It is generally thought that
the heterogeneous nanoscale morphology of hybrid BHJs is very critical towards high
performance of hybrid solar cells, as better percolation pathways would significantly
24
reduce charge recombination. This correlation has been directly exposed recently by Guo
et al. via transient absorption spectroscopy, showing that the electron transfer efficiency
for poly(2,6-(N-(1-octylnonyl)dithieno[3,2-b:20,30-d]pyrrole)-alt-4,7-(2,1,3-
benzothiadiazole)) PDBT: PbS NC BHJs can vary from 50% to 80% depending on the
nanoscale morphology.
73
However, it remains challenging to characterize and rationalize
the nanocomposite morphology experimentally. The morphology varies with several
experimental factors including choice of solvents, donor acceptor loading ratio, surface
ligands, deposition conditions, and post-deposition treatments. Most of these processing
conditions are optimized empirically. For instance, Sun et al. used 1,2,4-trichlorobenzene
(TCE) as a co-solvent for P3HT: CdSe BHJs, which allows a vertical phase segregation
in the blend, with a higher density of NCs near the cathode and a high polymer density
near the anode.
74
The resulting PCE was improved to 2.8%. Liu et al. applied the same
concept to PDTPBT: PbSxSe1-x hybrid BHJs by varying the loading ratio of NCs. With
the optimized phase segregation in the blend, they demonstrated a record PCE of 5.5%.
29

1.3.5. Optimizing Device Architecture
Much effort has been directed at minimizing the severe charge recombination at the
heterogeneous interfaces in hybrid BHJs, leading to current densities as high as ~18 mA
cm
–2
for state-of-the-art devices.
54
However, the overall performance is still much lower
than conventional OPVs. To this end, attention has also been drawn to the another
limiting factor, the absorption efficiency ( a), which is only 17.5%
30
for the PCPDTBT:
CdSe hybrid BHJs. This is because of the relatively narrow absorption band of the
photoactive materials by the use of low band gap polymers and NCs (vide infra), and
25
relatively thin active layers. Increasing the active layer thickness will increase the
absorption efficiency, however, the internal quantum efficiency will be compromised due
to the limited exciton diffusion length (5 20 nm)
1
and the poor charge transport
properties. One potential approach to enhancing the absorption efficiency without
introducing charge recombination, is to modulate the spatial distribution of optical field
inside the active layer. The idea behind this is that the optical intensity of the active layer
can be affected by interference of the incident light and the light reflected from the metal
electrode (e.g., Al). Therefore, by tuning the spatial position of the active layers through
varying the thickness of adjacent layers, or applying a multijunction device configuration,
the absorption efficiency of hybrid devices can be optimized. For example, it has been
demonstrated in OPVs that by inserting a transparent layer with an optimized thickness
between the active layer and electrode, the optical field intensity of the active layer can
be increased.
75, 76
Meanwhile, the transparent layer often serves as an electron/hole
blocking layer, thus preventing interfacial charge recombination. The same concept has
been applied to hybrid BHJ devices by Qian et al. in 2011.
77
The authors incorporated a
solution-processed ZnO NC layer between the active layer and the cathode, giving a
~70% enhancement in JSC for hybrid P3HT: CdSe BHJ devices. The improvement in JSC
was also revealed in the EQE, especially in the longer-wavelength region (Figure 1.9 a).
When CdSe NCs were blended with PCPDTBT, the optimized PCE of the hybrid devices
was enhanced to 3.5% with the insertion of ZnO layer, as compared to 2.3% without ZnO
layer. The improvement of the device efficiency was attributed to the elevated optical
electric field inside the active layer due to spatial redistribution (Figure 1.9 b).
26
Figure 1.9 (a) EQE spectra of hybrid P3HT: CdSe and P3HT: CdSe with a ZnO NC layer. (b)
Optical intensity profile for the device with the ZnO layer (  = 400 700 nm), showing that the
optical field intensity of BHJ is maximized. (c) Double-pass absorption percentage of single-
junction and tandem hybrid PH3T: CdSe solar cells. (d) I V curves of hybrid P3HT: CdSe single-
junction (47 nm active layer), tandem (47/52 nm), and triple-junction (47/47/56 nm) solar cells.
(a) and (b) are adapted with permission from [77].  2011 Royal Society of Chemistry. (c) and
(d) are adapted with permission from [78].  2014 American Chemical Society.

Another approach to improving the absorption efficiency is to build multi-
junction devices. This concept has been practiced in conventional OPVs.
79-81
For
example, Yang et al. showed that the light absorption in visible region can be increased
from 70% to 90% by a double-junction tandem OPV solar cells with two identical
subcells
82
as compared to a single-junction OPV. More importantly, the PCE was
improved from 8.1% in single junction OPV to 10.2% in tandem cells, stemming from
the two-fold increase in VOC over that of single-junction devices. We proved that the
same concept can be applied to hybrid BHJ solar cells. With nanocrystalline ZnO and
pH-neutral poly(3,4-ethylenedioxythiophene):poly(styrene sulfonic acid) (N-
PEDOT:PSS) as the interconnecting layers, tandem hybrid solar cells were successfully
fabricated with two identical subcells based on P3HT: CdSe BHJs.
78
The resulting
tandem cells present an increase of light absorption in the region of 480 600 nm from
27
60% to 80% compared to the single-junction hybrid devices (Figure 1.9 c). Meanwhile, it
gives a two-fold increase in VOC over that of single-junction hybrid devices, resulting in a
PCE of 2.7%. It is worth noting that the optical electric field of each active layer in the
tandem devices was optimized according to optical modeling using the transfer matrix
formalism, which further corroborated with device performances. As an extension of this
approach, triple-junction hybrid solar cells were also fabricated, giving a three-fold
increase of VOC (1.9 V) (Figure 1.9 d). These results clearly indicate that
tandem/multijunction device geometries are an effective approach to harvest more light
and improve PCE in hybrid BHJ solar cells. This may be further expanded to
multijunction hybrid devices containing nonequivalent subcells with complementary
absorption from both the selection of polymers and the composition and/or size of NCs,
offering additional parameter space for device optimization.

1.3.6. Harvesting NIR Light by Colloidal Lead Chalcogenide NCs
One of the main advantages of hybrid polymer: NC BHJ solar cells over fullerene-based
OPVs is the enhanced light absorption due to the NCs. State-of-the-art hybrid solar cells
are mostly based on CdSe NCs, in which absorption by the NC phase significantly
enhances photocurrent and EQE.
60, 65
However, the NIR part of the solar spectrum, which
contains 32% of the energy, is not harvested because of the intrinsic bulk band gap of
CdSe (Eg = 1.7 eV, corresponding to an absorption edge of ~730 nm). Meanwhile, there
has been a lack of conjugated polymers with relatively small band gaps because of the
synthetic limitations. These altogether have led to the usage of small band gap NCs, such
as lead chalcogenides (PbS and PbSe), in hybrid BHJ solar cells to extend the absorption
28
of devices into the NIR region. These lead chalcogenides possess band gaps in bulk of
0.41 (PbS) and 0.27 eV (PbSe). Combined with quantum confinement effects, lead
chalcogenides allow a broad-spectrum tunability in the NIR region. Additionally, lead
chalcogenides also possess relatively high dielectric constants (e.g., r ~ 17 for PbS)
83
and
the potential for MEG processes, both of which could be beneficial for hybrid devices. In
fact, Ginger et al. demonstrated via transient absorption spectroscopy that the excess
photon energy can be utilized without rapid phonon relaxation in lead chalcogenide NCs,
which can facilitate more efficient hole transfer from NCs to the polymer in a P3HT: PbS
hybrid system.
37
More recently, Guo et al. proved that the electron transfer process
between PDBT and PbS NCs happens in an ultrafast time scale of around 1 5 ps, with an
electron transfer efficiency of ~80%,
73
indicating that hybrid solar cells based on lead
chalcogenide NCs are indeed feasible in type II heterojunctions with suitable donor
polymers.
Although the concept of utilizing lead chalcogenide NCs to harvest NIR light is
promising, progress in hybrid solar cells based on polymer: PbE BHJs has been quite
slow.
84-88
In fact, it is not until recently that appreciable PCEs based on polymer:PbE
hybrid devices were obtained. In 2010, Ginger et al. explained that the poor performance
of hybrid polymer: PbS devices originates from a lack of photoinduced charge transfer at
the organic/inorganic interface.
89
This is due to the high lying valence band of PbS NCs,
preventing the formation of an energetically favorable type-Ⅱ heterojunction between the
polymer and the PbE NCs. For example, there is a small energy difference between the
valence band (5.0 5.2 eV) of PbS and the HOMO (5.1 eV) of P3HT, resulting in a small
driving force of hole transfer from PbS to the polymer. For this reason, the authors
29
utilized a new polymer, poly(2,3-didecyl-quinoxaline-5,8-diyl-alt-N-octyldithieno[3,2-
b:2’,3’-d]pyrrole) (PDTPQx), with a higher lying HOMO of 4.61 eV and confirmed a
significant photoinduced charge transfer in PDTPQx: PbS hybrid blends. After
exchanging the native oleate ligands with butylamine, hybrid PDTPQx: PbS devices
exhibited a PCE of 0.55%, which is 10 100 times larger than previously reported hybrid
devices based on lead chalcogenides. Prasad et al. further confirmed this argument by
introducing another low band gap polymer PDTPBT (Figure 1.1) as the electron donor.
67

After applying the postdeposition EDT treatment on the hybrid PDTPBT: PbS film, the
authors were able to achieve a high PCE of 3.8%, with a significant spectral contribution
from the NCs in the NIR region.
In addition to the search of suitable polymers, the surface chemistry of lead
chalcogenides NCs is an underdeveloped research area compared to cadmium
chalcogenides NCs. For instance, all of the high-performing hybrid polymer: PbS BHJ
devices mandate a postdeposition, thin film ligand treatment to remove native ligands
from the NC surface, which is a kinetically hindered process.
29,36,67,89-91
This is associated
with the fact that it can be especially difficult to obtain colloidally stable suspensions of
PbS NCs ligand exchanged with small ligands in solvents that are compatible with
polymer dissolution. While several successful examples of colloidal ligand exchange on
lead chalcogenide NCs using atomic halide ligands have been demonstrated, these
procedures require the ligand-exchanged NCs to be dispersed in polar solvents that are
orthogonal to the solubility of the semiconducting polymers.
43,92-96
Recently, we reported
a simple colloidal ligand exchange process for PbS NCs using iodide ligands (either NH4I
or PbI2), allowing them to be dispersed in 1,2-dichlorobenzene with the help of a small
30
amount of butylamine.
54
It is also worth noting that PbI2 passivation offers better
colloidal and air stability of PbS NCs. Moreover, it facilitates, for the first time, a one-
step deposition of hybrid Si-PCPDTBT: PbS BHJs to generate photovoltaic devices using
halide-passivated NC acceptors. The optimized hybrid solar cells display a broad spectral
response into the NIR, leading to a PCE of up to 4.8% for the PbI2-exchanged PbS NC
acceptors, which is much better compared to traditional solid-state EDT treatments
(Figure 1.10 a, b). Ultrafast transient absorption and time-resolved photoluminescence
spectroscopies suggest more efficient charge transfer and slower recombination of the
charge separated state in hybrid blends with PbI2 ligand treatment when compared to
hybrid films with NH4I surface ligands (Figure 1.10 c). More importantly, by selective
excitation of the PbS NC phase, we were able to exclusively probe the hole transfer
dynamics from the PbS NCs to Si-PCPDTBT by transient absorption spectroscopy,
which is a process otherwise often obscured by energy transfer.
90,97,98
A much longer hole
polymer polaron life time was observed in the hybrid Si-PCPDTBT: PbS BHJ with PbI2-
exchanged PbS acceptors, which was experimentally correlated to a more homogeneous
morphology of BHJ. In comparison, much shorter lifetime of hole polarons was observed
in the hybrid blend with NH4I-exchanged PbS NCs, indicating severe charge
recombination in the BHJs based on NH4I-exchanged PbS NCs. These results indicate
that the PbI2 ligand system is an attractive strategy for further hybrid solar cells with lead
chalcogenide acceptors.
31
Figure 1.10 Hybrid Si-PCPDTBT: PbS solar cells with different ligands on the surface of the
PbS NCs. (a) I V curves and (b) EQE spectra. (c) Femtosecond transient absorption spectra of
three films including a neat Si-PCPDTBT film, hybrid Si-PCPDTBT: PbS BHJ where PbS NCs
were capped with NH 4I ligands, and hybrid Si-PCPDTBT:PbS BHJ with PbS NCs exchanged
with PbI 2 ligands. Hybrid films were pumped at 920 nm and probed at 1200 nm, the neat polymer
film was pumped at 660 nm and probed at 1200 nm. Adapted with permission from [54].  2016
American Chemical Society.

1.4. Perspectives and Future Direction
Considering the combined benefits of both organic and inorganic phases, including
chemical tunability of energy gaps and interfacial states, solution processability, and
long-term stability, the performance of hybrid polymer:nanocrystal solar cells is still far
from initial expectations. However, the past two decades have witnessed a slow, yet
steady progress in both device performance and understanding of the mechanistic
processes of hybrid BHJ solar cells. For example, much effort has been devoted to
understanding the effects of NC size, shape, and surface ligands on the charge transfer
dynamics via both device characterization and spectroscopic evidence. In fact, it is
widely considered that these polymer: nanocrystal hybrid blends represent a very
interesting model system for studying the charge transfer and recombination processes in
photovoltaics.
30, 37, 99
This fundamental investigation will in turn provide direction for
further improving device performance of hybrid polymer:n anocrystal solar cells as well
as other types of photovoltaic devices.
Mechanistic investigation of hybrid BHJ solar cells points out that the poor
32
photovoltaic performance of devices stems from fast charge recombination, rather than
charge transfer at the D/A interface. In fact, charge transfer occurs at an ultrafast time
scale with a relatively high yield in both cadmium and lead based hybrid systems. Several
research groups recently ascribed the origin of the charge recombination to both the
spatial and electronic heterogeneity of organic and inorganic phases.
37,38,54,99
While the
spatial heterogeneity may be optimized through surface ligand engineering and
processing conditions, it remains a challenge to rationally control morphology.
Furthermore, tailoring the interfacial states of polymers and NCs strongly leans on the
development of surface chemistry of NCs, where it is particularly important to
understand the surface stoichiometry and trap states (and trap state density) of NCs.
Moreover, intrinsic electronic inhomogeneity may also be detrimental to photovoltaic
devices. For instance, the large asymmetry in dielectric constant, mobility and density of
electronic states between the polymer and NC phase likely yields imbalanced free
carriers, which subsequently contribute to trap-associated recombination. To this end,
along with spectroscopic investigations, future emphasis should be placed on the
following directions: (1) developing novel organic polymer donor materials and inorganic
colloidal NCs which possess compatible electronic properties (i.e., dielectric constant,
mobility, density of states), (2) designing surface ligands of NCs for a better coupling
between polymer-NC and NC-NC, (3) developing rational control over BHJ morphology.
All in all, the advantageous optoelectronic properties and flexibility of inorganic NCs
and continued progress in polymer chemistry suggest that hybrid polymer: nanocrystal
solar cells are a highly promising complement to the third generation photovoltaic
technology.
33
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42
Chapter 2. Tandem and Triple-Junction Polymer: Nanocrystal
Hybrid Solar Cells Consisting of Identical Subcells*
*Published in ACS Appl. Mater. Interfaces 2014, 6, 18306-18311.

2.1. Abstract
Tandem and triple-junction polymer: nanocrystal hybrid solar cells with identical subcells
based on P3HT: CdSe nanocrystal bulk heterojunctions (BHJs) are reported for the first
time showing two-fold and three-fold increases of open circuit voltage (VOC),
respectively, relative to the single junction cell. A combination of nanocrystalline ZnO
and pH-neutral PEDOT:PSS is used as the interconnecting layer and the thicknesses of
subcells are optimized with the guidance of optical simulations. As a result, the average
power conversion efficiency (PCE) exhibits a significant increase from 1.98% (VOC =
0.57 V) in single-junction devices to 2.7% (champion 3.1%, VOC = 1.28 V) in tandem
devices and 2.3% (VOC = 1.98 V) in triple-junction devices.
2.2. Introduction
Polymer: nanocrystal hybrid solar cells are receiving renewed attention because of the
potential they offer in combining beneficial properties of both the organic and inorganic
phases.
1−2
In place of the commonly employed fullerenes, inorganic semiconductor
nanocrystals can be used as alternate electron acceptors that take advantage of their
intrinsically higher electron mobilities, broad and tunable absorption at energies above
the band edge, and higher dielectric constants.
3−5
With this in mind, the power
43
conversion efficiency (PCE) of polymer: nanocrystal hybrid solar cells has improved to
>5%
6−7
over the past year. Proven strategies for increasing the PCE include changing the
nanocrystal size and shape,
8−10
ligand engineering
11−14,
and post-film deposition
treatments (e.g., ligand exchange via dip-coating,
6, 15
solvent-vapor annealing
16
). Despite
these performance gains, the PCEs of polymer: nanocrystal hybrid solar cells still lag
behind their all-organic counterparts. Earlier studies have shown that the main reasons
for poor device performance in hybrid solar cells is low absorption and limited charge
separation efficiency resulting from the limited exciton diffusion length.
17−19
One
potential solution for increasing absorption efficiency without increasing charge
recombination is applying a multijunction device configuration.
Parallel and series multijunction device configurations have been used in organic
photovoltaics (OPVs) to achieve enhanced short circuit current and open circuit voltage,
respectively, and have thusly achieved higher PCEs than their single-junction
counterparts.
20−24
When connected in series, open circuit voltages (VOC) that are the sum
of the VOC from individual subcells are theoretically obtained.
25
Despite the advantages
of multijunction over single-junction solar cells, there is only one example of a tandem
polymer: nanocrystal hybrid solar cell to the best of our knowledge. In that particular
case,
26
the PCE of the tandem solar cell connected in series ( P = 1.3%) was lower than
that of the corresponding single-junction device ( P = 1.8%), with only a 64% increase in
VOC relative to the single-junction cell being measured.
Herein, tandem and triple-junction polymer: nanocrystal hybrid solar cells with
identical poly(3-hexylthiophene) (P3HT): CdSe nanocrystal bulk heterojunction (BHJ)
subcells connected in series are reported. Interconnecting layers (ICLs) consisting of
44
nanocrystalline ZnO and pH-neutral poly(3,4-ethylenedioxythiophene):poly(styrene
sulfonic acid) (N-PEDOT:PSS) are used to successfully connect adjacent subcells.
Simulations are employed to optimize the thickness of each P3HT: CdSe BHJ by
maximizing the optical fields in each of the photoactive layers. As a result, tandem
polymer: nanocrystal hybrid solar cells giving a two-fold increase in VOC and up to ~50%
enhancement in PCE were obtained with an increase in absorption relative to single-
junction devices.

2.3. Results and Discussion
The as-synthesized CdSe nanocrystals were treated with a standard pyridine exchange to
remove the insulating native ligands in order to facilitate charge transfer between the
donor and acceptor phases of the BHJ.
11
The pyridine-exchanged CdSe nanocrystals
used in the tandem or triple-junction hybrid solar cells discussed herein possess a
diameter of 6.4 nm, as calculated from the max of the first exciton peak at 635 nm,
27

which is in agreement with TEM analysis. Thermogravimetric analysis of the pyridine-
exchanged nanocrystals shows a 5−6% mass loss up to 400 °C, which is indicative of loss
of organic ligands from the surface (Figure 2.1).
28
This compares favorably against the
as-synthesized CdSe nanocrystals (>15% mass loss up to 400 ° C), suggesting some
successful degree of ligand exchange.

45

Figure 2.1 (a) TGA data for the as-synthesized and pyridine-exchanged CdSe nanocrystals. (b)
TEM micrograph of the as-synthesized nanocrystals. Insert on the left is the size distribution and
insert on the right is a high-resolution image showing the lattice fringes of a single CdSe
nanocrystal. (c) Schematic and cross-section TEM micrograph of the tandem hybrid device
architecture: Glass/ITO/A-PEDOT:PSS/P3HT:CdSe(1)/ZnO/N-PEDOT:PSS/P3HT:CdSe(2)/ZnO
/Al.

2.3.1. Single-Junction Device Optimization
The performance of tandem/triple-junction OPVs is critically dependent on the
individual subcell active layer thicknesses. As a starting point, we sought to construct
single-junction P3HT: CdSe nanocrystal hybrid solar cells with an optimized active layer
thickness. The single-junction hybrid solar cell was fabricated with the device
configuration of ITO/PEDOT: PSS/P3HT: CdSe/ZnO/Al, with the P3HT: CdSe BHJ
active layer thickness varying between 42−62 nm. Figure 2.2a shows the I− V curves of
single-junction devices with different active layer thicknesses under 100 mW cm
−2
AM
1.5 G illumination, and the results are summarized in Table 2.1. The optimal active layer
thickness was found to be 47 nm, which gives an average PCE of 1.98(0.12)%, JSC =
6.95(0.33) mA cm
−2
, VOC = 0.57(0.01) V, and FF = 0.50(0.01). Active layer thicknesses
46
less than 47 nm lead to devices with low shunt resistance (RSH) and result in significant
decreases in FF when compared to the optimal thickness. The thicker active layers
exhibit larger RSH, but reduced JSC, likely resulting from enhanced charge recombination
owing to the relatively small exciton diffusion lengths in P3HT: CdSe BHJs.

Table 2.1 Photovoltaic devices parameters for single-junction P3HT: CdSe BHJ hybrid solar
cells with different active layer thicknesses.
a
Active layer
thickness (nm)
J SC (mA cm
−2
) V OC (V) FF PCE (%)
42 5.98  0.10 0.63  0.01 0.39  0.12 1.66  0.46
47 6.95  0.33 0.57  0.01 0.50  0.01 1.98  0.12
52 6.28  0.54 0.54  0.01 0.47  0.02 1.61  0.25
56 5.91  0.07 0.54  0.01 0.51  0.01 1.58  0.04
62 4.54  0.31 0.54  0.02 0.48  0.01 1.17  0.01
a
The device parameters were measured under AM 1.5G illumination at 1 sun. Average numbers
and the associated standard deviation(s) were determined over a total of 10−16 devices over four
separate substrates.

Figure 2.2 Single-junction P3HT:CdSe nanocrystal BHJ hybrid solar cell performance with
different active layer thicknesses. (a) I −V curves. (b) External quantum efficiency spectra. (c)
Internal quantum efficiency spectra.
47
To further investigate the active layer thickness dependence, the double-pass
absorption and external quantum efficiency (EQE) of the single-junction devices with
various thicknesses were measured and used to calculate the internal quantum efficiency
(IQE) by dividing the EQE spectra by the absorption percentage determined from
completed devices (Figures 2.2b and c). Photoresponse is observed out to ~700 nm,
which is the absorption edge of the 6.4 nm CdSe nanocrystals used in these devices, with
broadband IQE at higher energies coming from a combination of photoresponse from
both P3HT and the CdSe nanocrystals. It is noteworthy to point out that the peak IQE of
devices with a 47 nm (optimal) active layer thickness reaches 75%, and the values
decrease with further increasing thicknesses, indicating increased charge recombination
loss with thicker active layers. Thus, increasing the active layer thickness to harvest
more light is not an effective approach to increase the PCE of polymer: nanocrystal
hybrid solar cells, which could be potentially mediated by a tandem device architecture.

2.3.2. Tandem and Triple-Junction Devices
2.3.2.1 Interconnecting layer (ICL) optimization
The ICL is one of the most critical parameters to optimize in order to make an
operable tandem hybrid solar cell. Ideally, the ICL should work as a perfect transparent
ohmic contact (i.e., with no Schottky barrier). The combination of nanocrystalline ZnO
and neutral N-PEDOT:PSS has been previously demonstrated to be a robust ICL for
tandem/multijunction OPVs
29 30
connected in series. Here, we also observe successful
serial connection when this ICL is used between two adjacent P3HT: CdSe nanocrystal
BHJ subcells. It should be noted that the N-PEDOT:PSS is very important because
48
regular, acidic A-PEDOT:PSS (pH 1−2) will partially dissolve the underlying ZnO layer.
The concentration of ZnO and N-PEDOT:PSS were both optimized empirically based on
device performance. We find that 25 mg mL
−1
ZnO and 1: 2 dilution (v/v, diluted with
H2O) of N-PEDOT:PSS gives the best device performance. Those deposition conditions
result in 45 nm ZnO and 9 nm N-PEDOT:PSS layers as determined by spectroscopic
ellipsometry. Thicker N-PEDOT:PSS layers decrease the overall photocurrent due to the
partial light absorption by N-PEDOT:PSS, whereas thinner layers decrease the device
VOC, likely because of poor ohmic contact.
2.3.2.2 Tandem P3HT: CdSe nanocrystal subcell thickness optimization
Tandem hybrid solar cells with two identical subcells based on P3HT:CdSe
nanocrystal BHJs were fabricated as shown in Figure 2.1c, with the device architecture
ITO/A-PEDOT:PSS/P3HT: CdSe(1)/ZnO/N-PEDOT:PSS/P3HT: CdSe(2)/ZnO/Al. In
order to guide in the fabrication of the tandem solar cells, optical modeling using the
transfer matrix formalism
31
was performed on relevant devices architectures. Figure 2.3
depicts the output of simulated tandem device performance for the structure ITO/A-
PEDOT:PSS (27 nm)/P3HT: CdSe(1) (x nm)/ZnO (45 nm)/N-PEDOT:PSS (9 nm)/P3HT:
CdSe(2) (y nm)/ZnO (40 nm)/Al (100 nm). The simulations assume a constant IQE and
are meant to focus experimental efforts and are not an absolute predictor of device
performance. As a result of the optical cavity created by the reflective aluminum contact,
it is necessary to carefully control the thickness of the individual subcells so that they are
positioned in locations with maximized optical field intensity. The importance of this
behavior is reflected in the simulated photocurrent for cell 2 (Figure 2.3b), where, for a
given thickness for cell 1, the highest photocurrent is not achieved for the thickest cell 2,
49
but instead occurs at intermediate thickness where the cell encounters a maximum in
optical field strength. These results affirm that the ideal tandem device would contain
cells that are optimized for optical field intensity and not simply the largest thickness
attainable.
Additionally, because the two cells are connected in series where electrons
produced from the front subcell and holes produced from the rear subcell are recombined
in the interconnecting layer, the extractable photocurrent of the tandem hybrid solar cell
is limited by the individual cell with the smallest photocurrent. It is therefore important
to balance the photocurrent of the two cells to improve the total device efficiency. This is
illustrated in Figures 2.3c and 2.3d, which show the simulated photocurrent for the
tandem device and the photocurrent mismatch between the two cells, respectively. These
results highlight that the overall tandem performance is maximized at intermediate
thicknesses. Based on these simulations, the optimal tandem performance should be
achieved for individual subcell thicknesses around 50−60 nm.

Figure 2.3 (a) Simulated photocurrent from cell 1 in the tandem device as a function of active
layer thickness. (b) Simulated photocurrent from cell 2 in the tandem device as a function of
active layer thickness. (c) Photocurrent for the tandem device taken from the lowest photocurrent
50
of the two cells. (d) Current mismatch calculated from the difference in photocurrent between cell
1 and cell 2.

Experimentally, the thicknesses of both the front (1) and rear (2) active layers
were independently tuned from 47 nm to 62 nm. The details of the device performance
for these tandem hybrid solar cells are summarized in Table 2.2. The best performing
tandem hybrid solar cell has a front active layer thickness of 47 nm and a rear active layer
thickness of 52 nm, which is in close agreement with the results of the simulation (vide
supra). The champion tandem hybrid solar cells give JSC = 5.04 mA cm
−2
, VOC = 1.33 V,
FF = 0.47, and PCE = 3.12%. On average from 10 to 16 devices over four separate
substrates, the optimized tandem hybrid solar cells exhibit JSC = 4.61(0.39) mA cm
−2
, V OC
= 1.28(0.05) V, FF = 0.46(0.02), and PCE = 2.7(0.28)%. Surprisingly, the measured
average VOC (1.28 V) is 2.25 times of that of the average for an optimized single-junction
device. This is a result of higher shunt resistance in tandem devices, which in turn
decreases the dark current (Jo) in the tandem device. Figure 2.4 shows the dark current of
both single-junction device and tandem device. The dark current of tandem device at −1
V (0.05 mA cm
−2
) is 1/20th of that of the single-junction device (1 mA cm
−2
). Based on
the diode equation (𝑉 𝑂𝐶
=
𝑘 𝐵 𝑇 𝑞𝛽
(𝑙𝑛
𝑗 𝑆𝐶
𝑗 𝑜 + 1)), it is expected that lower dark current would
lead to higher VOC.
51

Figure 2.4 Dark I −V characteristic of single-junction and tandem P3HT: CdSe hybrid solar cells.
Table 2.2 Photovoltaic device parameters for tandem P3HT: CdSe BHJ hybrid solar cells with
different active layer (AL) thicknesses.
a

1
st
AL
Thickness
(nm)
2
nd
AL
Thickness
(nm)
J SC (mA cm
–2
) V OC (V) FF
PCE (%)
47
47 4.31  0.22 1.30  0.03 0.47  0.02
2.64  0.21
52 4.61  0.39 1.28  0.05 0.46  0.02
2.70  0.29
56 4.31  0.20 1.25  0.07 0.48  0.01
2.56  0.18
62 4.34  0.13 1.15  0.07 0.47  0.01
2.34  0.20

52
52 4.40  0.12 1.24  0.03 0.46  0.01
2.51  0.13
56 4.13  0.11 1.22  0.06 0.49  0.01
2.44  0.13
62 3.69  0.13 1.09  0.02 0.44  0.01
1.75  0.06

56
56 4.29  0.27 1.22  0.01 0.47  0.01
2.44  0.16
62 3.80  0.16 1.17  0.04 0.46  0.02
2.02  0.21

62 62 3.41  0.38 1.11  0.02 0.44  0.01
1.66  0.22

a
The device parameters were measured under AM 1.5G illumination at 1 sun. Average numbers
and the associated standard deviation(s) were determined over a total of 10 16 devices over four
52
separate substrates.

We also note that neither the VOC nor FF change significantly in the tandem
hybrid solar cells with different active layer thickness combinations; however, JSC does
change dramatically with the different active layer thickness combinations. The
maximum JSC of a tandem solar cell connected in series is limited by the subcell that
produces the smallest amount of photocurrent.
25
As such, individual subcells should be
designed to produce matched photocurrents, thereby maximizing the efficiency of
extracting photogenerated charge carriers. In our optimized tandem hybrid cell, we
observe an average JSC = 4.61(0.39) mA cm
−2
, which is 66% of the JSC measured for the
optimized single-junction device. When connected in series, 50% of the charge carriers
produced in each subcell are lost to recombination at the ICL.
25
Such a tandem device
should theoretically produce half the photocurrent of a single-junction device assuming
that both the absorption and IQE for the tandem and single-junction device are the same.
To understand the photocurrent enhancement observed in our system, the absorption of
both the optimized single and tandem hybrid P3HT: CdSe solar cells was measured in an
integrating sphere in reflectance mode (Figure 2.5). From these results, it is clear that the
absorption in the region of 480 nm to 600 nm in the tandem device is increased from 60%
to 80% relative to the single-junction device, thereby demonstrating that the tandem
configuration with two identical subcells does indeed improve light harvesting and lead
to higher photocurrent.

53
Figure 2.5 Double-pass absorption percentage of single-junction and tandem P3HT: CdSe hybrid
solar cells.

2.3.2.3 Triple-junction hybrid P3HT: CdSe solar cells
As a general proof of principle, triple-junction tandem hybrid solar cells were also
fabricated with three identical P3HT: CdSe nanocrystal BHJ active layers, and the
resulting devices exhibited a more than three-fold increase in VOC (VOC = 1.91(0.20) V)
relative to the single-junction device. The I − V curves of the optimized champion hybrid
solar cells with single-junction, double-junction, and triple-junction device geometries
are given in Figure 2.6, and the device parameters are summarized in Table 2.3. The
thicknesses of the active layers in the triple-junction hybrid solar cells were 47, 47, and
56 nm (from front to rear relative to incident irradiation). The average PCE of the triple-
junction devices ( P = 2.3%) is not higher than the double-junction devices because the
thicknesses of each subcell were not optimized to give the best current; however, the
open circuit potential was shown to be additive from those of the three subcells as
expected for a triple-junction device connected in series. Therefore, the combination of
54
nanocrystalline ZnO and pH-neutral PEDOT:PSS is a robust ICL for serial multijunction
hybrid solar cells.

Figure 2.6 I −V curves of P3HT: CdSe single-junction (47 nm active layer), tandem (47/52 nm)
and triple-junction (47/47/56 nm) hybrid solar cells.

Table 2.3 Photovoltaic devices details for optimized single-junction, tandem, and triple-
junction hybrid solar cells based on a P3HT: CdSe BHJ.
a
Devices J SC (mA cm
−2
) V OC (V) FF PCE (%)
Single device
b
6.95  0.33 0.57  0.01 0.50  0.01 1.98  0.12
Tandem device
c
4.61  0.39 1.28  0.05 0.46  0.02 2.70  0.29
Triple device
d
2.80  0.14 1.91  0.20 0.42  0.02 2.26  0.28
Champion
tandem device
5.04 1.33 0.47 3.12
a
The device parameters were measured under AM 1.5G illumination at 1 sun. Average numbers
and the associated standard deviation(s) were determined over a total of 10−16 devices over four
separate substrates.
b
Active layer thickness is 47 nm.
c
Active layers thicknesses are 47 nm
(front) and 52 nm (rear).
d
Active layers thicknesses are 47 nm (front), 47 nm (middle) and 56 nm
(rear).

55
2.4. Experimental
2.4.1. Materials
CdCO3 (99.998% metal basis, "Puratronic" grade, Alfa Aesar), selenium (200 mesh
powder, 99.999% metal basis, Alfa Aesar), tri-n-octylphosphine oxide (TOPO, 98%, Alfa
Aesar), tri-n-octylphosphine (TOP, ≥ 97%, Strem), stearic acid (95%, Sigma-Aldrich),
pyridine (≥ 99.0%, "GRACS" grade, EMD) were all used as received. P3HT (≥ 96% rr,
65−75 kDa MW, 1.7−1.9 PDI) was purchased from Reike Metals. Acidic A-
PEDOT:PSS (Clevios PH 500, pH 1−2, percent 1−1.5) was purchased from Heraeus.
Neutral N-PEDOT:PSS (high-conductivity grade, pH 5−7, percent 1.1) was purchased
from Sigma-Aldrich.

2.4.2. Synthesis and Ligand Exchange of CdSe Nanocrystals
The synthesis is based on literature methods.
27
In a typical synthesis, CdCO3 (3.50 g, 20.3
mmol), stearic acid (30.0 g), and tri-n-octylphosphine oxide (TOPO, 30.0 g) were stirred
at 100 ° C under flowing nitrogen for ~1.5 h. Then the temperature was increased to
360 ° C and held under static nitrogen for 1 h to completely dissolve the solids. With rapid
stirring, TOP/TOPSe (selenium (2.30 g, 29.1 mmol) previously dissolved in TOP (30 mL,
67 mmol) under nitrogen) was injected (~ 30 s). Exactly 2 min after the end of injection,
the flask was removed from the heating bath and the reaction was quenched via air-
cooling.
The reaction product was split between six 45-mL centrifuge tubes. EtOH (20 mL
each tube) was then added to the centrifuge tubes, and the mixture was centrifuged (6000
rpm, 2 min) and the supernatant was removed. This washing procedure was repeated four
56
times to remove most of the free ligands on the nanocrystal surface. Small amounts of
agglomerates were carefully removed and the final dispersion was made in 70 mL of
toluene with 1 mL of oleic acid to prevent agglomeration. The product was stored in the
dark at 10 ° C. This synthesis procedure gives CdSe nanocrystals with a first exciton peak
near 635 nm, which corresponds to a nanocrystal diameter of 6.4 nm.
Pyridine exchange of CdSe nanocrystals: A dispersion of the as-synthesized CdSe
nanocrystals in toluene (10 mL, 400 mg CdSe) was added to 40 mL of pyridine in a 100-
mL round-bottomed flask fitted with a small water-cooled condenser. The system was
vigorously purged with nitrogen for 10 min, and then introduced into a 115 ° C sand bath
and heated for 12 h. After cooling, the dispersion was divided between six 45-mL
centrifuge tubes. Pentane (40 mL in each tube) was added, the mixture was centrifuged
(6000 rpm, 2 min), and the supernatant was discarded. Pyridine was immediately added
to disperse the solid. The final dispersion was made in 20 mL of pyridine, which easily
passed through a 0.45 m PTFE syringe filter to give the final suspension, which was
stored in the dark at room temperature.

2.4.3. Characterization
Absorption spectra were acquired on a Perkin-Elmer Lamba 950 spectrophotometer
equipped with a 150 mm integrating sphere, using a quartz cuvette for liquid samples or a
borosilicate glass microscope slide substrate for films. Thermogravimetric analysis
(TGA) measurements were made on a TA Instruments TGA Q50 instrument, using
sample sizes between 5−15 mg in an alumina crucible under a flowing nitrogen
atmosphere. TGA samples were prepared by drying the colloid under flowing nitrogen at
57
80 C for up to 90 min, then lightly crushing with a spatula. Film thicknesses were
determined using a J. A. Woollam variable angle spectroscopic ellipsometer equipped
with a 150 W Xe arc lamp. TEM images were obtained on a JEOL JEM-2100F
microscope at an operating voltage of 200 kV, equipped with a Gatan Orius CCD camera.

2.4.4. Single-Junction Hybrid Solar Cell Fabrication
All devices were fabricated and tested in air. Aluminum shot (Al; Alfa Aesar, 99.999%)
was purchased and used as received. Patterned ITO-coated glass substrates (10  cm
−2
,
Thin Film Devices, Inc.) were sequentially cleaned by sonication in tetrachloroethylene,
acetone, and isopropanol followed by 30 min of UV-ozone treatment. A 27 nm
(determined by ellipsometry) layer of A-PEDOT:PSS (Clevios PH 500, pH 1−2, percent
1−1.5, 1: 1 v/v diluted by H2O, filtered through a 0.45 m cellulose acetate syringe filter)
was spun-cast (4000 rpm, 40 s) onto the clean ITO and heated at 120 C for 30 min under
vacuum (~5 kPa). A P3HT solution of 15 mg mL
−1
was prepared in 1,2-dichlorobenzene
by dissolution under mild heating (40−50 C) followed by filtering through a 0.45 m
PTFE syringe filter. The pyridine exchanged CdSe nanocrystals were dispersed in a
mixed solvent of 90% 1,2-dichlorobenzene (DCB) and 10% pyridine by volume, and
filtered through a 0.45 m PTFE syringe filter. To make the final blend solution, the
CdSe nanocrystals (in 90% DCB, 10% pyridine) were then mixed for several hours under
mild heating with the filtered P3HT solution to a final concentration of 3: 24 mg mL
−1

(P3HT: CdSe). The active layer was spun-cast at different speeds (2500, 2000, 1500,
1000, and 700 rpm, all for 50 s) onto the dried PEDOT: PSS layer to get active layer
thicknesses of 42, 47, 52, 56, and 62 nm, respectively. After drying in a dark nitrogen
58
cabinet for 20−25 min, ZnO nanocrystals dispersed in ethanol (20 mg mL
−1
) were spun-
cast on the active layer (4000 rpm, 40 s) to produce a 40 nm thick layer. The devices
were then annealed at 150 C under flowing nitrogen for 10 min, followed by loading
into a high vacuum (~2 Torr) thermal deposition chamber (Angstrom Engineering) for
deposition of 100-nm thick Al cathodes through a shadow mask at a rate of 2 Å s
−1
.
Device active areas were 4.3 mm
2
as measured by pixel mapping through a CCD-
equipped optical microscope.

2.4.5. Tandem/Triple-Junction Hybrid Solar Cell Fabrication
The cleaning procedure for ITO substrates was identical for all devices fabricated. A 27
nm layer of A-PEDOT: PSS (Clevios PH 500, pH 1−2, percent 1−1.5, 1: 1 v/v diluted by
H2O, filtered through a 0.45 m cellulose acetate syringe filter) was spun-cast (4000 rpm,
40 s) onto the clean ITO and heated at 120 C for 30 min under vacuum (~5 kPa). The
first active layer was spun-cast at different speeds (2000, 1500, 1000, and 700 rpm, all for
50 s) onto dried A-PEDOT: PSS layer to get an active layer thickness of 47, 52, 56, and
62 nm, respectively. After drying in a dark nitrogen cabinet for 20−25 min, ZnO
nanocrystals dispersed in ethanol (25 mg mL
−1
) were spun-cast on the active layer (4000
rpm, 40 s) to produce a 45 nm layer. Then the device was annealed at 150 C under
flowing nitrogen for 10 min. A 9 nm layer of N-PEDOT: PSS (pH 5−7, percent 1.1, 1:2
v/v diluted by H2O, filtered through a 0.45 m cellulose acetate syringe filter) was spun-
cast (6000 rpm, 40 s) onto the ZnO layer and heated at 120 C for 30 min under vacuum
(~5 kPa). The second active layer was spun-cast at different speeds (2000, 1500, 1000,
and 700 rpm, all for 50 s) onto the dried N-PEDOT: PSS layer. Then, the second ZnO
59
layer (20 mg mL
−1
) was spun-cast on the second active layer (4000 rpm, 40 s) to produce
a 40 nm thick layer, and the device was annealed at 150 C under flowing nitrogen for 10
min. To make triple-junction hybrid solar cells, the third N-PEDOT: PSS, active layer,
and ZnO layers are processed in the same way as the second junction of the tandem
hybrid solar cell. Finally, the devices were loaded into a high vacuum (~2 Torr) thermal
deposition chamber (Angstrom Engineering) for deposition of 100-nm thick Al cathodes
through a shadow mask at a rate of 2 Å s
−1
.

2.4.6. Device Characterization
Current-density dependence on applied test voltage measurements were performed under
ambient conditions using a Keithley 2400 SourceMeter (sensitivity = 100 pA) in the dark
and under ASTM G173−03 spectral mismatch corrected 1000 W m
−2
white light
illumination from an AM 1.5G filtered 450 W Xenon arc lamp (Newport Oriel).
Chopped and filtered monochromatic light (250 Hz, 10 nm FWHM) from a Cornerstone
260 1/4 M double grating monochromator (Newport 74125) was used in conjunction with
an EG&G 7220 lock-in amplifier to perform all spectral responsivity measurements.

2.5. Conclusions
We present the first successful example of tandem and triple-junction hybrid polymer:
nanocrystal solar cells with identical subcells based on P3HT: CdSe nanocrystal BHJs.
The tandem hybrid solar cells give a two-fold increase in VOC over that of single-junction
devices, and light absorption in the region of 480 nm to 600 nm is increased from 60% to
60
80% compared to the single-junction hybrid solar cells. To further prove the robustness
of this device architecture, triple-junction hybrid solar cells were also fabricated and were
shown to produce three times the VOC of single-junction devices. The PCE exhibits a
substantial increase from 2.0% in single-junction devices to 2.7% (champion 3.1%) in
tandem devices and 2.3% in triple-junction devices. The results clearly indicate that
tandem/multijunction device geometries are an effective approach to harvest more light
and improve PCE in hybrid BHJ solar cells. This may be further extended to
multijunction hybrid devices containing nonequivalent subcells with complementary
absorption from both the selection of polymers and the composition and/or size of
quantum dots, which offers additional parameter space for device optimization.

2.6. References
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(2) McGehee, M. D., Nanostructured Organic–Inorganic Hybrid Solar Cells. MRS
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(13) Radychev, N.; Lokteva, I.; Witt, F.; Kolny-Olesiak, J.; Borchert, H.; Parisi, J. r.,
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(14) Greaney, M. J.; Brutchey, R. L. Ligand Engineering in Hybrid
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(15) Seo, J.; Cho, M. J.; Lee, D.; Cartwright, A. N.; Prasad, P. N., Efficient
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62

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(19) Gao, F.; Li, Z.; Wang, J.; Rao, A.; Howard, I. A.; Abrusci, A.; Massip, S.;
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(22) You, J.; Dou, L.; Yoshimura, K.; Kato, T.; Ohya, K.; Moriarty, T.; Emery, K.;
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(25) Ameri, T.; Dennler, G.; Lungenschmied, C.; Brabec, C. J., Organic tandem solar
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(26) Kaltenhauser, V.; Rath, T.; Edler, M.; Reichmann, A.; Trimmel, G., Exploring
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64
Chapter 3. Iodide-Passivated Colloidal PbS Nanocrystals leading to Highly
Efficient Polymer: Nanocrystal Hybrid Solar Cells*
*Published in Chem. Mater. 2016, 28, 1897–1906.

3.1. Abstract
Current state-of-the-art hybrid polymer: lead chalcogenide nanocrystal solar cells require
post-deposition, thin film chemical treatments to remove insulating organic ligands from
the nanocrystal surface, which is a kinetically hindered process. This is compounded by
the fact that it can be especially difficult to obtain colloidally stable suspensions of PbS
nanocrystals ligand exchanged with small ligands, and many atomic ligands require
dispersion in solvents that are incompatible with polymer solubility. Herein, we report a
novel one-step colloidal ligand-exchange process for PbS nanocrystals employing lead
iodide (PbI2) or ammonium iodide (NH4I) as surface ligands along with n-butylamine that
allow the ligand exchanged nanocrystals to be suspended in solvents compatible with
polymer dissolution. While ligand exchange is shown to be near quantitative for both
iodide sources, when compared to NH4I-exchanged PbS nanocrystals, the PbI2-
exchanged PbS nanocrystals not only exhibit better colloidal stability but also display
superior photovoltaic performance. When the PbI2-passivated PbS nanocrystals are
combined with the donor polymer poly[2,6-(4,4´ -bis(2-ethylhexyl)dithieno[3,2-b:2´ 3´ -
d]silole)-alt-4,7-(2,1,3-benzothiadiazole) (Si-PCPDTBT), the optimized hybrid solar cells
give a broad spectral response into the NIR leading to a power conversion efficiency
(PCE) of 4.8% under AM 1.5G illumination. Time-resolved photoluminescence and
transient absorption spectroscopies suggest that excitonic processes between the PbS
65
nanocrystals and Si-PCPDTBT are more favorable when PbS nanocrystals are ligand
exchanged with PbI2, leading to superior device performance.

3.2. Introduction
Organic polymer: fullerene bulk heterojunction (BHJ) solar cells offer the opportunity for
low-cost device fabrication as a result of the solution processibility of both active
components, in addition to having the added benefit of employing a thin absorber layer
that translates to low materials costs.
1,2
While organic photovoltaics have now achieved
power conversion efficiencies in excess of ~10%,
3 5
optimization of the organic donor
phase may be nearing its limit to further increase device performance. This has led to
increased emphasis on exploring new acceptor types.
6,7
Semiconductor nanocrystals
have been explored as alternatives to the more well-established fullerene acceptors, and
they possess several attributes that make them attractive in this regard: (i) size- and
composition-tunable absorption from the visible to near infrared (NIR), (ii) intrinsically
higher electron mobilities, (iii) the potential for multiple exciton generation (MEG), and
(iv) and higher dielectric constants to help overcome the strong exciton binding energy of
organic materials.
8 10
To date, hybrid solar cells based on polymer: CdSe nanocrystal
BHJs have been extensively explored and optimized with respect to ligand
engineering,
11 16
size and shape effects,
17 20
charge transfer dynamics,
21 24
and device
architectures,
25 28
leading to power conversion efficiencies (PCEs) of 4 5%. However,
the best performing polymer: CdSe nanocrystal BHJs only absorb visible light out to
wavelengths that are limited by the donor polymer, which prevents them from harvesting
66
light into the NIR part of the solar spectrum. For example, CdSe nanocrystals are
intrinsically limited by the bulk band gap (Eg = 1.7 eV) to an absorption edge of ~730
nm. Moreover, the dielectric constant of bulk CdSe ( r ~6
29
) is not significantly different
from that of the organic polymer phase ( r ~34
30
), whereas lower band gap
semiconductors generally possess high dielectric constants (e.g., r ~17 for PbS
31
).
Because of these attributes, lead chalcogenide (PbE, where E = S, Se) nanocrystals are
being increasingly investigated as electron acceptors in hybrid polymer: nanocrystal BHJ
solar cells.
32 35

Although the photovoltaic performance of hybrid polymer: PbE nanocrystal BHJ
solar cells are promising, much less is known about the effects of ligand engineering for
the lead chalcogenides on device performance when compared to CdSe acceptors.
36,37

All of the high-performing, state-of-the-art hybrid polymer: PbE nanocrystal BHJ devices
to date mandate a post-deposition, thin film ligand treatment to remove insulating native
ligands from the nanocrystal surface.
32 37
Such ligand exchange processes are extremely
critical to provide the interparticle electronic coupling that is required for efficient
excitonic processes to occur (i.e., charge transfer, charge separation). Common bidentate
ligands (e.g., 1,2-ethanedithiol (EDT), 3-mercaptoproprionic acid) and atomic halide
ligands (e.g., iodide) have been applied via thin film ligand exchange;
36
however, thin
film ligand exchange has several inherent drawbacks: (i) ligand exchange can be impeded
by slow and/or incomplete solid-state ligand diffusion through the BHJ film, (ii)
reduction of the film volume through exchange of small ligands for large ones can lead to
film cracking, and (iii) the dip-washing process used for thin film ligand exchange can
67
lead to a low ligand exchange efficiency (yield) and may also be incompatible with large
scale solution processing.
38

With this in mind, a quantitative colloidal ligand exchange prior to BHJ
deposition would be beneficial, but it still remains a challenge to generate a stable
suspension of colloidal lead chalcogenide nanocrystals exchanged with small organic
ligands because of their tendency to agglomerate,
37
etch, and/or oxidize after ligand
exchange.
39
For example, simple alkylamines (e.g., n-butylamine or BA) typically
require 48 72 h to achieve incomplete ligand exchange, and lead to surface etching and
poor colloidal stability for PbS nanocrystals in air.
39
Colloidal ligand exchange of PbS
nanocrystals with arenethiolate ligands still required a post-deposition, thin film ligand
treatment with 3-mercaptopropionic acid to achieve PCEs of 3% in hybrid poly(3-
hexylthiophene) (P3HT): PbS nanocrystal BHJ solar cells.
37
On the other hand, atomic
halide ligands have proven successful in achieving high efficiency colloidal quantum dot
solar cells,
40 42
and in addition to allowing for efficient interparticle coupling, have been
shown to effectively passivate surface trap states in PbS nanocrystals.
43
While there are
several successful examples of colloidal ligand exchange of lead chalcogenide
nanocrystals with atomic halide ligands,
44 47
these procedures require the ligand-
exchanged nanocrystals be dispersed in polar solvents that are orthogonal to the solubility
of semiconducting polymers. Therefore, we sought to exploit the beneficial attributes of
halide passivated PbS nanocrystals that are arrived at through a colloidal ligand
exchange, but are dispersible in solvents that are compatible with the donor polymers.
Herein, we report a simple colloidal ligand exchange process for PbS nanocrystals
using iodide ligands (either NH4I or PbI2) that allows them to be dispersed in 1,2-
68
dichlorobenzene with the help of a small amount of n-butylamine. This facilitates, for the
first time, a one-step deposition of hybrid poly[2,6-(4,4´ -bis(2-ethylhexyl)dithieno[3,2-
b:2´ 3´ -d]silole)-alt-4,7-(2,1,3-benzothiadiazole) (Si-PCPDTBT): PbS nanocrystal BHJs
to generate photovoltaic devices with halide-passivated nanocrystal acceptors. The
optimized hybrid solar cells exhibit a broad spectral response into the NIR, leading to a
PCE up to 4.8% under AM 1.5G illumination for the PbI2-exchanged PbS nanocrystal
acceptors. This device performance is among the best for hybrid polymer: nanocrystal
solar cells, and is the highest for a one-step BHJ deposition performed without any post-
deposition ligand exchanges. Results from ultrafast transient absorption (TA) and time-
resolved photoluminescence (PL) spectroscopies indicate more efficient charge transfer
and slower recombination of the charge separated state in hybrid films with PbI2 ligand
treatment when compared to hybrid films with NH4I surface ligands.

3.3. Results and Discussion
3.3.1. Colloidal Iodide Ligand Exchange
It has been previously shown that PbS nanocrystals with a 1.3 eV band gap result in the
best charge separation efficiency in Si-PCPDTBT: PbS nanocrystal BHJs.
35
Therefore,
PbS nanocrystals with an average diameter of 3.2 nm (determined from TEM, Figure 3.1)
and a band gap of 1.3 eV were synthesized from a modified hot-injection approach.
48

The as-synthesized PbS nanocrystals are initially capped with oleate native ligands,
which provide excellent colloidal stability in nonpolar organic solvents.
69

Figure 3.1 TEM micrograph of the as-synthesized PbS nanocrystals. Inset on the right is a high-
resolution image showing the lattice fringes of a single PbS nanocrystal.

We developed a methanol-assisted procedure for exchanging PbS nanocrystals
with inorganic iodides, such as NH4I or PbI2. In general, the inorganic ligands are
dissolved in methanol (for NH4I), or a mixture of methanol and DMF (1: 2 vol/vol for
PbI2), followed by addition of a PbS nanocrystal suspension. Methanol acts as an anti-
solvent and can partially remove the surface X-type ligands (i.e., oleate),
51
leading to
nanocrystal precipitation, and the inorganic ligands can simultaneously passivate the
nanocrystal surface and enable dispersibility in polar solvents. The solvent dispersibility
of these iodide-exchanged PbS nanocrystals can be easily changed into medium polarity
1,2-dichlorobenzene with the addition of n-butylamine (40% by volume for NH4I and
20% by volume for PbI2), which enables solution compatibility with Si-PCPDTBT. The
integrity of the PbS nanocrystals is retained after this iodide ligand exchange, as
confirmed by solution UV-vis-NIR spectroscopy (Figure 3.1a). The absorption spectra of
70
the iodide-exchanged PbS nanocrystals show the persistence of sharp excitonic features,
albeit slightly red-shifted (~37 meV) when compared to as-prepared PbS nanocrystals
with oleate ligands. This slight red shift may be attributed to a change in surface dipole
moments or solvent dielectric constants.
42
Thermogravimetric analysis (TGA) data was
used to measure the mass loss upon heating from room temperature to 400 C under
flowing nitrogen (Figure 3.2b). TGA provides information on the total mass of ligands
through their decomposition/volatilization temperature relative to that of the inorganic
nanocrystal core.
52
It is observed that after ligand exchange, the amount of organic
content is reduced. This is evidenced by a lower mass loss percentage at 400 ° C, and is
indicative of the efficient displacement of long-chain oleate native ligands during the
ligand exchange process.
53 55
The organic ligand content decreases from 30% for oleate-
capped PbS nanocrystals to 4.5 and 3.6% after NH4I and PbI2 ligand exchange,
respectively. The TGA data are supported by FT-IR analysis of the PbS nanocrystals,
which exhibit a significant decrease in  C H) stretching intensity between 3000 2800
cm
–1
(Figure 3.2c), suggesting a quantitative or near-quantitative removal of the native
oleate ligands.
1
H NMR spectroscopy further confirms the completeness of ligand
exchange, as evidenced by a dramatic decrease in the integrated intensity in the aliphatic
region between 0–5 ppm (Figure 3.3). The presence of surface iodide is confirmed by
high-resolution X-ray photoelectron spectroscopy (XPS) for both the NH4I- and PbI2-
exchanged PbS nanocrystals, with binding energies of 619 and 631 eV being observed for
the I 3d5/2 and 3d3/2 peaks, respectively (Figure 3.2d). The atom% surface iodide
concentration, quantified from the XPS survey spectrum, is 30% relative to sulfur for
PbI2 treatment and 44% relative to sulfur for NH4I. These data collectively suggest that
71
the efficacy of iodide exchange is excellent when compared to traditional BA ligand
exchange, which only partially removes the native oleate ligands, as evidenced by TGA,
FT-IR spectroscopy, and
1
H NMR (Figure 3.3). For instance, ~15 wt% of oleate ligands
remain on the surface of the PbS nanocrystals after a 48 h BA treatment. Additionally,
the hypsochromic shift of excitonic features in the UV-vis-NIR absorption spectrum
suggests that ligand exchange with BA results in significant etching of the PbS
nanocrystal surface.
39
Figure 3.2 Colloidal ligand exchange characterization of the PbS nanocrystals. (a) vis-NIR
absorption spectra, (b) TGA traces, and (c) FT-IR spectra of the PbS nanocrystals with different
surface ligands (OA = oleate, BA = n-butylamine, NH 4I, and PbI 2). (d) Normalized high-
resolution I 3d XPS spectra for NH 4I- and PbI 2-exchanged PbS nanocrystals.

72

Figure 3.3 Normalized
1
H NMR spectra of oleate-, BA-, NH 4I-, and PbI 2-passivated PbS
nanocrystals after digestion in aqua regia and extraction with d 6-benzene.

The iodide-exchanged PbS nanocrystals exhibit very different colloidal stabilities
depending on the halide source; that is, PbI2-exchanged PbS nanocrystals are colloidally
stable for up to two months in 1,2-dichlorobenzene, while brown precipitates are
observed in suspensions of NH4I-exchanged PbS nanocrystals within 24 h after ligand
exchange. Such differences in surface chemistry might be attributed to the introduction
of Pb
2+
counterion resulting from: (i) the surface passivation of sulfur sites with Z-type
ligand binding,
56
and/or (ii) the formation of [PbI3]
–
X-type ligands
46
(as evidenced by
negative-ion electrospray ionization mass spectrometry of PbI2 in DMF, Figure 3.4a). An
observed increase in the Pb to S ratio after PbI2 ligand exchange from 1.38 to 1.48, based
on inductively coupled plasma optical emission spectroscopy, further supports these
hypotheses. It is also worth noting that the PbI2-exchanged PbS nanocrystals display
outstanding air stability. PbS nanocrystals are known to be sensitive to air, forming
PbSOx after exposure to the ambient environment in both solution and thin films.
39
As a
73
result, a blue shift of the first excitonic peak is observed after exposure to air. However,
the PbI2-capped PbS nanocrystals display exactly the same excitonic features after being
exposed to ambient air conditions for one month (Figure 3.4b), demonstrating that the
PbS nanocrystals ligand exchanged with PbI2 offer exceptional colloidal and air
stabilities.

Figure 3.4 (a) ESI-MS spectrum of PbI 2 dissolved in DMF solution (100 g mL
–1
). The peak at
126.9 m/z indicates the presence of I
–
, while the peaks at 586.7 and 588.7 m/z indicate the
presence of PbI 3
–
. (b) Normalized vis-NIR absorption spectra of PbS nanocrystals with different
surface ligands before and after 30 d aging under ambient conditions. Top spectra are for PbI 2-
exchanged PbS nanocrystals, while the bottom spectra are for oleate-passivated PbS nanocrystals.

3.3.2. Hybrid Solar Cells
74
Figure 3.5 (a) Schematic illustration of hybrid solar cells based on ITO/PEDOT:PSS/Si-
PCPDTBT: PbS nanocrystal/ZnO/Al device structures. (b) Normalized absorbance spectra of
PbS nanocrystal film (black), of Si-PCPDTBT film (red) and a hybrid Si-PCPDTBT: PbS BHJ
film (blue).

A Si-PCPDTBT: PbS nanocrystal BHJ absorber layer is expected to possess a
type-II band alignment between the donor and acceptor phases (Figure 3.5a).
35
This
active layer was then incorporated into hybrid solar cells with the general device structure
ITO/PEDOT: PSS (35 nm)/Si-PCPDTBT: PbS nanocrystal BHJ/ZnO (40 nm)/Al (100
nm) via spin-casting blends of Si-PCPDTBT with PbS nanocrystals dispersed in 1,2-
dichlorobenzene that were filtered through a 0.45 µ m filter. The devices were optimized
with respect to the polymer/nanocrystal loading ratio, followed by the active layer
thickness, and finally the iodide ligand concentration used during ligand exchange
(Figure 3.6 and Table 3.1). It was determined that the Si-PCPDTBT: PbS nanocrystal
BHJ that performs best has a loading ratio of 30:2 mg mL
–1
Si-PCPDTBT: PbS, an active
layer thickness of 48 nm, and a PbI2 concentration of 0.02 M (18 mL) for exchanging 100
mg of PbS nanocrystals in 2 mL of toluene (Table 3.2).

75
Figure 3.6 Photovoltaic device PCE for hybrid Si-PCPDTBT: PbS BHJ solar cells with (a)
different PbI 2 ligand concentrations (thicknesses of BHJ are all 59 nm), and (b) different BHJ
film thicknesses.

Table 3.1 Photovoltaic device parameters for hybrid Si-PCPDTBT: PbS nanocrystal BHJ solar
cells with different active layer (AL) thicknesses.
a
AL Thickness (nm) J SC (mA cm
–2
)
b
V OC (V) FF PCE (%)
42 15.17  0.38 0.49  0.01 0.49  0.01 3.64  0.26
48 16.60  1.55 0.48  0.01 0.50  0.04 3.98  0.80
59 13.12  0.38 0.47  0.01 0.41  0.04 2.46  0.58
83 7.75  1.00 0.46  0.01 0.35  0.03 1.25  0.29

a
The device parameters were measured under AM 1.5G illumination at 1 sun. Average
numbers and the associated standard deviation(s) were determined over a total of 10−16
devices over 4 separate substrates.
b
Short-circuit current densities (J SC) were corrected by
mismatch factor.

The highest performing hybrid solar cells under optimized conditions come from
PbI2-capped PbS nanocrystals. The champion device gave JSC = 18.2 mA cm
–2
, VOC =
0.48 V, FF = 0.55, and PCE = 4.78%. This champion device performance is among the
best for hybrid solar cells to date. Average device parameters were calculated from up to
16 devices over four separate substrates for hybrid Si-PCPDTBT: PbS nanocrystal BHJ
solar cells where the acceptor had been ligand exchanged with PbI2, giving JSC = 16.6
(1.6) mA cm
–2
, VOC = 0.48 (0.01) V, FF = 0.50 (0.04), and PCE = 3.98 (0.80)%. Strong
photoresponse from both Si-PCPDTBT and the PbI2-exchanged PbS nanocrystals can be
observed in the external quantum efficiency (EQE) (Figure 3.7), which reveals a broad
spectral response from the UV into the NIR (i.e., 3001100 nm). This BHJ combination
gives a peak EQE (at 650 nm) of 60%, with a NIR EQE (at 980 nm) of 10% originating
from the PbS nanocrystal acceptors. In comparison, hybrid solar cells based on BHJs of
Si-PCPDTBT: PbS nanocrystals ligand exchanged with NH4I gave considerably lower
photocurrent density and fill factor (FF), with average device parameters of JSC = 6.4
76
(0.9) mA cm
–2
, VOC = 0.43 (0.04) V, FF = 0.31 (0.01), and PCE = 0.87 (0.18)%. In this
device, a lower peak EQE (at 500 nm) of 46% was observed, in addition to a much
weaker EQE response of 2% in the NIR at 980 nm. Moreover, benchmark control
devices were also fabricated with two alternate ligand treatments on the PbS nanocrystal
acceptors: (i) colloidal BA ligand exchange and (ii) colloidal BA ligand exchange
followed by a thin film ligand exchange with EDT after the BHJ was spin cast. These
control systems also exhibit lower performance with respect to the BHJ utilizing the PbI2-
exchanged PbS nanocrystals, possessing average PCEs of 1.27 (0.19)% and 2.80 (0.60)%
for BA- and EDT-exchanged PbS nanocrystals, respectively (Table 3.2). Of particular
interest to us, however, are the apparent differences in EQE response and short circuit
current density between the two different iodide ligand exchanges (i.e., PbI2 vs NH4I),
which may suggest differences in charge transfer and separation in the two systems. To
further investigate the origin of the performance differences between these two iodide
ligand treatments, time-resolved PL and ultrafast TA spectroscopic studies were
conducted.

77
Figure 3.7 Performance of hybrid Si-PCPDTBT: PbS nanocrystal BHJ solar cell devices with
different ligands on the surface of the PbS nanocrystals. (a) I V curves and (b) external quantum
efficiency spectra.

Table 3.2 Photovoltaic device raw data for hybrid Si-PCPDTBT: PbS nanocrystal BHJ solar cells
with different ligands.
a
Ligands
J SC, Raw
(mA cm
–2
)
M
b
J SC, Corr
(mA cm
–2
)
c

V OC (V) FF PCE (%)
OA 3.60  0.29 0.98 3.68  0.29 0.33  0.10 0.29  0.02 0.36  0.12
BA 5.35  0.90 0.86 6.22  1.05 0.55  0.02 0.37  0.02 1.27  0.19
BA-EDT 8.96  1.20 0.80 11.20  1.50 0.54  0.01 0.46  0.05 2.80  0.60
NH 4I 5.80  0.80 0.90 6.40  0.90 0.43  0.04 0.31  0.01 0.87  0.18
PbI 2 12.8  1.20 0.77 16.60  1.55 0.48  0.01 0.50  0.04 3.98  0.80
Champion-
PbI 2
14.0 0.77 18.2 0.48 0.55 4.78

a
The device parameters were measured under AM 1.5G illumination at 1 sun. Average
numbers and the associated standard deviation(s) were determined over a total of 10−16
devices over 4 separate substrates.
b
Mismatch factor was calculated from M =
∫ 𝐸 𝑅 ( )𝑆 𝑅 ( )𝑑 

2

1
∫ 𝐸 𝑅 ( )𝑆 𝑇 ( )𝑑 

2

1
∫ 𝐸 𝑆 ( )𝑆 𝑇 ( )𝑑 

2

1
∫ 𝐸 𝑆 ( )𝑆 𝑅 ( )𝑑 

2

1
, where E R( ) is the reference spectral irradiance; E S( ) is the
source spectral irradiance; S R( ) is the spectral responsivity; and S T( ) is the spectral
responsivity of the test cell, each as a function of wavelength ( ). Spectral responsivities
S( ) for tested devices were calculated based on the external quantum efficiency (EQE)
values, according to equation S( ) =
𝑞 
ℎ𝑐 𝐸𝑄𝐸 ( ).
c
J SC, Corr =
𝐽 𝑆𝐶 ,𝑅𝑎𝑤
𝑀 ).

3.3.3. Charge Separation Dynamics at the Hybrid Interface
To probe the dynamics of charge transfer or energy transfer between the Si-PCPDTBT
donor and the PbS nanocrystal acceptors that have been ligand exchanged with PbI2 and
NH4I, time-resolved PL quenching studies were performed. Time-resolved PL spectra
were collected for a neat Si-PCPDTBT thin film and Si-PCPDTBT: nanocrystal BHJ thin
films containing PbI2- and NH4I-exchanged PbS (and prepared under identical processing
conditions to the working devices). A significant reduction of PL lifetime for Si-
PCPDTBT at 700 nm is observed for hybrid BHJs containing both NH4I- and PbI2-
exchanged PbS acceptors when the films were excited at 500 nm (Figure 3.8). At this
78
excitation wavelength both the polymer and nanocrystals are excited. Energy transfer
from the excited state of the polymer can lead to an excited state of the nanocrystals,
while electron transfer from the excited polymer to nanocrystals can result in a charge
separated state. The same charge separated state can also be formed by hole transfer
from the excited nanocrystal to the polymer. The measured PL lifetime of Si-PCPDTBT
(  = 1.25 ns) is longer than a previous literature report of 250 ps.
35
The extended PL
lifetime observed here might result from extra precautions taken to avoid exposure of the
polymer thin film to air. For the hybrid BHJ films, a bi-exponential decay dynamic is
observed for the Si-PCPDTBT: nanocrystal BHJ with NH4I-exchanged PbS (Figure 3.8).
The lifetimes obtained by fitting the decay traces to a bi-exponential decay function were
0.387 ns (9%) and 0.026 ns (91%), while the excited state for the Si-PCPDTBT:
nanocrystal BHJ with PbI2-exchange PbS decays faster than our instrument response
function (<22 ps). These results suggest that the PbI2 ligand exchange enables a more
efficient charge transfer or energy transfer than NH4I exchange based on time-resolved
PL measurements, which correlates with JSC measured in the hybrid devices.

79
Figure 3.8 Time-resolved PL traces of the pristine Si-PCPDTBT polymer and Si-PCPDTBT:
nanocrystal BHJs employing NH 4I- and PbI 2-exchanged PbS acceptors (  ex = 500 nm;  em = 700
nm). The instrument response function (IRF) is shown in black.

Ultrafast TA spectroscopy was used in a complementary way to gain deeper
insight into the charge transfer dynamics. Previous TA studies for hybrid polymer: PbS
systems were based on excitation at wavelengths where both the polymer and PbS
nanocrystals possess considerable absorbance.
35,57
As a result, both energy transfer and
charge transfer between the polymer and PbS nanocrystals can occur simultaneously, as
in our time-resolved PL quenching experiments, thereby complicating the interpretation
of the spectral dynamics. In our TA measurements, the PbS nanocrystals were
selectively excited by pumping at 920 nm, where Si-PCPDTBT does not absorb, and the
polaron dynamics of Si-PCPDTBT were probed at 1200 nm. This wavelength was
chosen based on an assignment to a polymer hole polaron at 1200 nm for its carbon-
bridged analogue C-PCPDTBT in a previous literature report.
57
To confirm this
assignment for the particular Si-PCPDTBT polymer used in this study, we carried out
chemical doping experiments by oxidizing the polymer with SbCl5.
50
Chemical treatment
of Si-PCPDTBT with SbCl5 results in oxidation of the polymer, similar to the charge
separated state we expect to observe in the TA experiments after hole transfer from the
PbS nanocrystals to the polymer upon excitation at 920 nm. The steady-state absorption
spectra of the chemically oxidized Si-PCPDTBT displays a reduction in intensity of the
400 and 660 nm bands, concomitant with the appearance of a broad positive band around
1200 nm (Figure 3.9), which is very similar to what has been previously observed for C-
PCPDTBT.
50

80

Figure 3.9 (a) Steady-state absorption spectra of the Si-PCPDTBT thin film before doping
(black) and after a 3 min dip in a 20 ppm solution of SbCl 5 in acetonitrile (red), shown together
with the difference spectrum ( A = A doped  A undoped, given in blue). (b) Transient absorption
spectra of neat Si-PCPDTBT film (at 0.1 ps, pumped at 660 nm) overlapping with chemically
doped steady-state spectrum of the oxidized Si-PCPDTBT film. Steady-state absorption spectra
were taken in an integrating sphere and thus account for scattering and reflections.

The steady-state absorption spectrum of the chemically oxidized polymer does not
exactly overlap in the NIR range with TA spectra obtained for the neat Si-PCPDTBT
polymer (excited at 660 nm, Figure 3.9b) and the hybrid BHJ films (excited at 920 nm)
because of (i) a (positive) photoinduced absorption at 1400 nm for the singlet exciton of
the polymer also contributes to the TA spectra of the neat polymer and, (ii) to the spectra
of hybrid films, a bleach at 920 nm from the PbS nanocrystals, and, specifically for NH4I-
exchanged nanocrystals, a strong stimulated emission band at 1100 nm are clearly
contributing negative-going signals (Figure 3.10).
50

81

Figure 3.10 Transient absorption spectra (at 4 ps, pumped at 920 nm) overlapping with the
absorption and emission spectra of the PbS nanocrystals. (a) Hybrid Si-PCPDTBT: nanocrystal
BHJ film with PbI 2-exchanged PbS, and (b) hybrid Si-PCPDTBT: nanocrystal BHJ film with
NH 4I-exchanged PbS.

Quantitative deconvolution of the TA spectra in the NIR will be conducted in the
future; however, the hole polaron does contribute to the positive 1200 nm signal of the
TA spectra at the 0.1 ps decay time scale (Figure 3.9b) based on the overlapping spectral
feature from chemical oxidation. Therefore, we probed the positive band at 1200 nm to
gain insight into the dynamics of the polymer hole polaron. The formation of polymer
hole polarons is direct evidence of a hole transfer process from the ligand-exchanged PbS
nanocrystals to Si-PCPDTBT, without any interference from energy transfer.
82

Figure 3.11 Transient absorption spectra of a neat Si-PCPDTBT film (pumped at 660 nm) (a),
and hybrid Si-PCPDTBT: nanocrystal BHJ films with (b) PbI 2- and (c) NH 4I-exchanged PbS
acceptors (pumped at 920 nm) at different time delays between the pump and the probe pulse.

Figure 3.11 shows the full transient absorption spectra of a neat Si-PCPDTBT
film (pumped at 660 nm) (a), and hybrid Si-PCPDTBT: nanocrystal BHJ films with (b)
PbI2- and (c) NH4I-exchanged PbS acceptors (pumped at 920 nm) at different time delays
between the pump and the probe pulse. The evolution of the band at 1200 nm with pump-
probe delay for the corresponding films is given in Figure 3.12. In all films, the polaron
83
signal appears within the instrument response (≤ 120 fs), suggesting good coupling
between the PbS nanocrystals and the Si-PCPDTBT polymer for either ligand
treatment.
21
While the 1200 nm induced absorption in the neat polymer shows a bi-
exponential decay with lifetimes of  = 0.6 ps and 53 ps, the induced absorption decays
much faster in the Si-PCPDTBT: nanocrystal BHJ with NH4I-exchanged PbS acceptors,
but persists for a considerably longer time in the BHJ film with PbI2-exchanged PbS
acceptors. The 1200 nm kinetics in the hybrid BHJ films are clearly affected by the
presence of bleach from the PbS nanocrystals as well as the induced absorption from the
polaron in the same spectral window (Figure 3.10). In the case of the NH4I ligand
exchange, a stimulated emission signal is also clearly observed in the complex TA trace
for the hybrid BHJ film. This suggests that a fraction of the NH4I-exchanged PbS
nanocrystals do not charge separate in the hybrid BHJ film, perhaps because of large
domains of the NH4I-exchanged PbS acceptors (vide infra) or poor coupling; however,
the fast polaron rise time seen for the NH4I-exchanged nanocrystals that do charge
separate suggest a heterogeneous model. A detailed kinetic analysis of the various
signals seen in the broadband TA experiment using global analysis is underway and will
be the subject of future work. However, it is evident from Figure 3.12 that the lifetime of
the polaron is much longer in hybrid BHJ films with PbI2-exchanged PbS nanocrystal
acceptors as compared to the neat polymer film, and clearly different from the hybrid
BHJ films with NH4I-exchanged PbS nanocrystal acceptors. Such a long-lived hole
polaron on the polymer in hybrid BHJ films with PbI2-exchanged PbS nanocrystal
acceptors would thus allow the holes to better percolate toward the electrode.
84

Figure 3.12 Femtosecond TA spectra of three films as labeled. Hybrid BHJ films were
pumped at 920 nm and probed at 1200 nm; the neat Si-PCPDTBT film was pumped at
660 nm and probed at 1200 nm.

The differences in the lifetimes of the polaron in hybrid BHJs with NH4I- and
PbI2-exchanged PbS nanocrystals may be attributed to secondary hole transfer processes
(e.g., nanocrystal surface traps or reverse hole transfer from polaron to nanocrystals). To
test these hypotheses, PL spectra of PbS nanocrystal suspensions and the morphologies of
the hybrid BHJ films were investigated. It was determined that the concentration-
normalized PL intensity of the NH4I- and PbI2-exchanged PbS nanocrystal suspensions
were both lower than the as-prepared PbS nanocrystals with oleate ligands (Figure 3.13),
suggesting that both iodide exchanges result in more surface trap states than the oleate-
passivated PbS nanocrystals. Such mid-gap surface trap states might account for the
lower VOC of the devices with iodide-exchanged PbS nanocrystal acceptors relative to the
benchmark devices utilizing BA and BA-EDT ligand treatments. Additionally, both the
NH4I- and PbI2-exchanged PbS nanocrystal suspensions exhibit similar concentration
normalized PL intensities, suggesting qualitatively similar densities of surface trap states.
85
Therefore, we cannot specifically ascribe the significant differences in excited state
lifetimes to surface trap states on the nanocrystal acceptors.

Figure 3.13 Concentration-normalized steady-state PL spectra of PbS nanocrystal suspensions in
1,2-dichlorobenzene with different surface ligands: oleate (OA), NH 4I, and PbI 2 (  ex = 900 nm).
The feature around 1140 nm is due to an instrument artifact from a grating change in the
fluorimeter.

However, a considerable difference in the BHJ morphology between the two
ligand treatments is observed, however, by atomic force microscopy (AFM) and
transmission electron microscopy (TEM). The hybrid Si-PCPDTBT: nanocrystal BHJ
with NH4I-exchanged PbS possesses a rough film surface (rms roughness of 72 nm) with
a large degree of nanocrystal aggregates, while that with PbI2-exchanged PbS exhibits a
much smoother film morphology (rms roughness of 16 nm) (Figures 3.14 a,b). Such
morphology variation can also be observed in the corresponding TEM images (Figures
3.14 c,d). The nanocrystal aggregates observed in the hybrid BHJs with NH 4I-exchanged
PbS nanocrystals might provide a pathway for secondary reverse hole transfer from
86

Figure 3.14 (a b) AFM topological image of a hybrid Si-PCPDTBT: nanocrystal BHJ film with
NH 4I- and PbI 2-exchanged PbS, respectively. Image was obtained using 10  10 μm window.
(c d) TEM image of a hybrid Si-PCPDTBT: nanocrystal BHJ film with NH 4I- and PbI 2-
exchanged PbS nanocrystals (d = 3.2 nm). White circles in (c) indicate nanocrystal aggregates
with domain sizes ranging from 10 to 40 nm. While in (d), individual nanocrystals can be seen
embedded in the polymer matrix.

polaron to nanocrystals,
57
leading to a lower lifetime of the polaron, and/or result in a
failure to charge separate. These processes can explain the enhanced stimulated emission
observed in the TA spectra from the NH4I-exchanged PbS nanocrystals in the hybrid film
(Figure 3.10b), which is caused by an enhanced excited state population of the
nanocrystals compared to PbI2-exchanged nanocrystals. The tendency of NH4I-
exchanged PbS nanocrystals to aggregate in the hybrid BHJ films is also consistent with
their behavior in the solution phase, where NH4I-exchanged PbS nanocrystals aggregate
87
much more easily than the PbI2-exchanged nanocrystals. These observations are
consistent with the EQE spectra of the hybrid BHJ devices in the NIR region, where no
distinct NIR exciton peak from the PbS nanocrystal acceptors is generally observed,
except for the hybrid Si-PCPDTBT: nanocrystal BHJ with PbI2-exchanged PbS
nanocrystals. Therefore, nanocrystal aggregation in the hybrid Si-PCPDTBT:
nanocrystal BHJ with NH4I-exchanged PbS may account for the poor charge separation
dynamics.

3.4. Experimental
3.4.1. Materials
Lead oxide (PbO, 99.99%), lead iodide (PbI2, 99.9985%), and ammonium iodide (NH4I,
99.999%), n-butylamine (99%) were purchased from Alfa Aesar. Bis(trimethylsilyl)
sulfide ((TMS)2S, >95.0%) was purchased from TCI America. Oleic acid (90%), 1,2-
ethanedithiol (99%) and 1-octadecene (90%) were purchased from Sigma-Aldrich.
PEDOT:PSS (Clevios PH 500, pH 1−2, percent 1−1.5) was purchased from Heraeus. Si-
PCPDTBT (40,500 MW, 2.2 PDI) was purchased from 1-Material. All chemicals are
used as received without further purification.
3.4.2. Synthesis of PbS Nanocrystals
The synthesis is based on literature methods.
58
In a typical synthesis, 0.92 g (4.0 mmol)
of PbO, 4.0 g (13 mmol) of oleic acid and 20 g of octadecene were added into a 100 mL
three-neck flask with a water-cooled condenser. The mixture was then heated and stirred
88
continuously under vacuum at 90 ˚C overnight, forming a clear and colorless lead oleate
solution. 420 L of (TMS)2S (2.00 mmol) was combined with a 10 g dry octadecene and
loaded into a syringe. Caution: (TMS)2S is extremely reactive and it can react with
moisture in air rapidly producing H2S gas! The (TMS)2S in octadecene solution was
rapidly injected into the lead oleate solution under nitrogen. After 3 s, the solution turned
dark, indicating the nucleation of PbS nanocrystals. The heating mantle was immediately
removed after injection (but with stirring), and the solution was allowed to cool to room
temperature naturally. PbS nanocrystals were then isolated by flocculation with excess
acetone and redispersion in hexanes. The final suspension was made in 20 mL of toluene
with a concentration of 50 mg mL
–1
.
3.4.3. Ligand Exchange of PbS Nanocrystals
3.4.3.1. Ligand Exchange with Lead Iodide (PbI2)
A 0.02 M solution of PbI2 in a mixture of DMF (12 mL) and methanol (6 mL) was
prepared first. The suspension of as-prepared PbS nanocrystals (2 mL, 50 mg mL
–1
) was
then added into the PbI2 solution and shaken for 1 2 min, leading to an immediate
precipitation of the PbS nanocrystals. After centrifugation, the PbS nanocrystals were
isolated from the clear supernatant, followed by redispersing them in a mixture of 1,2-
dichlorobenzene (1 mL) and BA (0.2 mL). The ligand-exchanged PbS nanocrystals were
then filtered through a 0.45 m PTFE syringe filter for device fabrication. The PbI2-
exchanged PbS nanocrystals are stable up to several months without any sign of
agglomeration or surface etching.
89
3.4.3.2. Ligand Exchange with Ammonium Iodide (NH4I)
First, a 0.2 M solution of NH4I in methanol (8 mL) was prepared. The suspension of as-
prepared PbS nanocrystals (2 mL, 50 mg mL
–1
) was then added into the NH4I solution,
leading to an immediate precipitation of the PbS nanocrystals. After centrifugation, the
PbS nanocrystals were isolated from the clear supernatant, followed by redispersing them
in a mixture of 1,2-dichlorobenzene (1 mL) and BA (0.4 mL). The ligand-exchanged
PbS nanocrystals were then filtered through a 0.45 m PTFE syringe filter for device
fabrication. The NH4I-exchanged PbS nanocrystals are not quite colloidally stable, and
brown solids will precipitate within 48 h.
3.4.3.3. Ligand Exchange with N-Butylamine (BA)
A dispersion of as-prepared PbS nanocrystals (2 mL, 50 mg mL
–1
) was first precipitated
by 20 mL of acetone. After the supernatant was removed, BA (8 mL) was then added to
redisperse the PbS nanocrystals. The solution was stored in the dark under a nitrogen
atmosphere for 2 d. The PbS nanocrystals were then precipitated with excess acetone (30
mL) and redispersed in 1 mL of DCB. The ligand-exchanged PbS nanocrystals were then
filtered through a 0.45 m PTFE syringe filter for device fabrication. The BA-exchanged
PbS nanocrystals are not quite colloidally stable, and brown solids will precipitate within
24 h.
3.4.4. Characterization
UV-vis-NIR absorption spectra were acquired on a Perkin-Elmer Lamba 950
spectrophotometer equipped with a 150 mm integrating sphere, using a quartz cuvette for
90
liquid samples or a borosilicate glass microscope slide substrate for films.
Thermogravimetric analysis (TGA) measurements were made on a TA Instruments TGA
Q50 instrument, using sample sizes between 5-15 mg in an alumina crucible under a
flowing nitrogen atmosphere. TGA samples were prepared by drying the colloid under
flowing nitrogen at 80 C for up to 90 min, then lightly crushing with a spatula prior to
analysis. FT-IR spectra were acquired from pressed pellets on a Bruker Vertex 80.
Pressed pellets were made of dried nanocrystals (3 mg) and an internal standard, Prussian
blue Fe4[Fe(CN)6]3, in a dry KBr matrix (100 mg) in order to gain semi-quantitative
information.
1
H NMR spectra were collected on a Varian 500 spectrometer (500 MHz in
1
H) with chemical shifts represented in units of ppm. All spectra are normalized relative
to the residual benzene solvent peak at 7.16 ppm for the purpose of enabling semi-
quantitative comparison between different samples. NMR sample preparation and
analyses were conducted according to previously published procedures for nanocrystal
ligand analysis.
49
Briefly, samples were prepared by digesting dried nanocrystal solids
(~50 mg) in half-concentrated aqua regia (7 mL), and then the organics were extracted
using d6-benzene (2 mL). The solvent was then dried with MgSO4 and filtered before
1
H
NMR analysis. 32 scans were taken for each sample, and the data are presented as
averages of those scans. XPS spectra were obtained using a Kratos Axis Ultra X-ray
photoelectron spectrometer with an analyzer lens in hybrid mode. High resolution scans
were performed using a monochromatic aluminum anode with an operating current of 5
mA and voltage of 10 kV using a step size of 0.1 eV, a pass energy of 20 eV, and a
pressure range between 1 3  10
–8
torr. The binding energies for all spectra were
referenced to the C1s core level at 284.6 eV. Inductively coupled plasma optical
91
emission spectroscopy (ICP-OES) was performed on each nanocrystal sample for
analysis of lead and sulfur using a Thermo Scientific Icap 7000 series ICP-OES. Film
thicknesses were determined using a J. A. Woollam variable angle spectroscopic
ellipsometer equipped with a 150 W Xe arc lamp. TEM images of BHJ films supported
on copper (Ted Pella, Inc.) were obtained on a JEOL JEM-2100F microscope at an
operating voltage of 200 kV, equipped with a Gatan Orius CCD camera.
3.4.5. Photoluminescence Lifetime Studies
For time correlated single photon counting (TCSPC) measurements, the samples were
spun cast onto borosilicate glass in a glovebox with optical densities between 0.1 0.2 at
the excitation wavelength of 500 nm. To avoid any oxidative damage, the film also had
an additional glass window placed on the top surface and the outer edges were sealed
with epoxy under a nitrogen atmosphere. Lifetime measurements were carried out using
the output of a Coherent RegA 9050 regenerative amplifier operating at 250 kHz. The
amplified 800 nm pulse was then used to pump an optical parametric amplifier, Coherent
OPA 9450, which produced the 500 nm excitation beam used in the experiment. The
lifetime data were measured using pump pulses of power 0.5, 0.67, and 1.7 mW and with
a spot size of 250 μm at the sample, resulting in excitation fluences of 4.0, 5.4 and 14 μJ
cm
–2
for the neat Si-PCPDTBT and hybrid Si-PCPDTBT: nanocrystal BHJ films with
NH4I- and PbI2-exchanged PbS acceptors, respectively. PL lifetimes were measured by
detecting the emission at 700 nm for the neat polymer and hybrid films. The film
samples were not moved for the duration of the experiment and remained static with
respect to the excitation beam. TCSPC measurements were performed using a R3809U-
50 Hamamatsu PMT with a B&H SPC-630 module (time resolution of 22 ps). The
92
monochromator grating was blazed at 600 nm with 1200 g mm
–1
and a slit width of 1.2
mm was used, giving a spectral bandpass of 4.17 nm. The lifetime measurements were
limited by the response of the detector (22 ps), which was longer than the pulse width of
the excitation beam from the OPA (<200 fs).
3.4.6. Transient Absorption Studies
Broadband femtosecond TA experiments were carried out using the output of a Coherent
Legend Ti:sapphire amplifier (1 kHz, 3.5 mJ, 35 fs). The 900 1200 nm supercontinuum
probe pulses were generated in a sapphire window driven by the 800 nm amplifier
fundamental, polarized perpendicular to the pump pulses and detected with an InGaAs
Hamamatsu G9213-256S photodiode array after being dispersed by an Oriel MS1271
spectrograph. The 920 nm pump was generated by doubling 1840 nm generated using
~10% of the amplifier’s output to seed a type II collinear OPA (Spectra Physics OPA-
800C). To minimize probe dispersion, a pair of off-axis aluminum parabolic mirrors was
used to collimate the probe and focus it into the sample, while a CaF2 lens focused the
pump. Samples were translated perpendicular to the path of the pump and probe to
prevent photodamage. TA spectra were measured with a pump energy 0.14 μJ and a spot
size of 120 μm resulting in a fluence of 1300 μJ cm
–2
. Spectra were acquired by
averaging each time slice over 500 laser shots, and delay traces were scanned 5 times.
3.4.7. Chemical Doping Studies
In order to characterize the steady state polaron spectrum of Si-PCPDTBT, thin films
were spun cast on quartz slides from a 2 mg mL
–1
solution in 1,2-dichlorobenzene. Si-
PCPDTBT films were subsequently oxidized by dipping them in a 20 ppm solution of
93
SbCl5 in acetonitrile.
49
The state-state absorption spectra were acquired on a Perkin-
Elmer Lamba 950 spectrophotometer equipped with a 150 mm integrating sphere to
account for scattering and reflections. The difference spectra were acquired by
subtracting the spectrum of the oxidized polymer from that of the neat polymer.
3.4.8. Hybrid Solar Cell Fabrication
All devices were fabricated in air and tested under nitrogen atmosphere. Aluminum shot
(Al; Alfa Aesar, 99.999%) was purchased and used as received. Patterned ITO-coated
glass substrates (10  cm
−2
, Thin Film Devices, Inc.) were sequentially cleaned by
sonication in tetrachloroethylene, acetone, and isopropanol, followed by 20 min of UV-
ozone treatment. A 35 nm hole transporting layer of PEDOT:PSS (Clevios PH 500, pH
1−2, percent 1−1.5, filtered through a 0.45 m cellulose acetate syringe filter) was spun-
cast (4000 rpm, 40 s) onto the clean ITO and heated at 120 ˚C for 30 min under vacuum
(~5 kPa). A Si-PCPDTBT solution of 15 mg mL
−1
was prepared in 1,2-dichlorobenzene
by dissolution under mild heating (30 40 ˚C) for 1 h, followed by filtering through a 0.45
m PTFE syringe filter. The ligand-exchanged PbS nanocrystals (in 1,2-dichlorobenzene
and BA) were then mixed for 1 h at room temperature with the filtered Si-PCPDTBT
solution to a final concentration of 2: 30 mg mL
–1
(Si-PCPDTBT: PbS). This solution
was spun cast at 1500 rpm for 50 s onto the dried PEDOT:PSS layer to get a 48 nm thick
active layer, and it was aged in the dark under flowing nitrogen for 20 25 min. For the
EDT-exchanged BHJ as control devices, the BHJ was then dipped into a 1% (vol/vol)
solution of EDT in dry acetonitrile for 1 min, followed by washing with dry acetonitrile
twice to remove residue organic ligands. ZnO nanocrystals (synthesized from a sol-gel
method)
59
dispersed in ethanol (20 mg mL
−1
) were spun cast on the active layer (4000
94
rpm, 40 s) to produce a 40 nm thick electron transport layer. The devices were then
annealed at 130 ˚C under flowing nitrogen for 10 min, followed by loading into a high
vacuum (~2 Torr) thermal deposition chamber (Angstrom Engineering) for deposition
of 100-nm thick Al cathodes through a shadow mask at a rate of 2 Å s
−1
. Device active
areas were 4.4 mm
2
as measured by pixel mapping through a CCD-equipped optical
microscope.
3.4.9. Device Characterization
Current-density dependence on applied test voltage measurements were performed under
ambient conditions using a Keithley 2400 SourceMeter (sensitivity = 100 pA) in the dark
and under ASTM G173−03 spectral mismatch corrected 1000 W m
−2
white light
illumination from an AM 1.5G filtered 450 W Xenon arc lamp (Newport Oriel).
Chopped and filtered monochromatic light (250 Hz, 10 nm FWHM) from a Cornerstone
260 1/4 M double grating monochromator (Newport 74125) was used in conjunction with
an EG&G 7220 lock-in amplifier to perform all spectral responsivity measurements. An
unfiltered silicon photodiode (400−1100 nm) calibrated by NREL was used as a reference
for spectral responsivity measurements.

3.5. Conclusions
We have demonstrated a facile and near-quantitative colloidal ligand exchange process
for PbS nanocrystals based iodide ligands from PbI2 and NH4I. Hybrid solar cells based
on BHJs of Si-PCPDTBT and PbS nanocrystals were fabricated from the direct solution
deposition of polymer and nanocrystal mixtures without further solid-state ligand
95
exchange. Si-PCPDTBT: nanocrystal BHJ devices with PbI2-exchanged PbS acceptors
achieved a PCE of 4.8%; on the other hand, NH4I-exchanged PbS nanocrystals exhibit
significantly lower photovoltaic performance when blended with Si-PCPDTBT under the
same conditions. Time-resolved PL spectroscopy indicates a more efficient energy
transfer or charge transfer process occurs when PbI2-exchanged PbS nanocrystals are
blended with Si-PCPDTBT, which is consistent with the measured short circuit current
density and integrated EQE for these devices. To further elucidate the exciton dynamics
at the hybrid interface, we were able to selectively probe the hole transfer dynamics from
the PbS nanocrystals to Si-PCPDTBT by ultrafast TA spectroscopy via the selective
excitation of the nanocrystal phase. A much longer hole polaron lifetime is qualitatively
observed in the hybrid Si-PCPDTBT: nanocrystal BHJ with PbI2-exchanged PbS
acceptors. This taken together with excellent colloidal and air stability, facile one-step
device fabrication, and superior photovoltaic performance render the PbI2 ligand system
attractive for future hybrid solar cells applications with lead chalcogenide acceptors.

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101
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102
Chapter 4. Exposing the Dynamics and Energetics of the N-Heterocyclic
Carbene Nanocrystal Interface *
*Published in J. Am. Chem. Soc. 2016, 138, 14844–14847.

4.1. Abstract
N-heterocyclic carbenes (NHCs) are becoming increasingly popular ligand frameworks
for nanocrystal surfaces; however, as of yet, the nature of the NHC-nanocrystal interface
remains unexplored across different material types. Here, we report a facile synthetic
route to NHC-stabilized metal and metal chalcogenide nanocrystals. It was observed that
NHC-Ag nanocrystals are colloidally stable, but much less so than the corresponding
NHC-Ag2E analogs. Comprehensive NMR studies suggest a dynamic NHC-nanocrystal
interface for both NHC-Ag and Ag2S; however, DFT calculations reveal a much stronger
binding affinity of the NHC ligands to Ag2S compared to Ag nanocrystals, which
explains the superior colloidal stability of the metal chalcogenides. This offers new
insight into the surface chemistry of neutral, L-type carbenes in colloidal nanocrystal
chemistry.

4.2. Introduction
N-heterocyclic carbenes (NHCs) are an important class of neutral, L-type ligands owing
to their structural diversity, chemical stability, and electron-rich σ-donating ability that
leads to stable bonds with a wide swath of elements in the periodic table.
1
As a direct
result of these characteristics, NHCs have been gradually gaining interest as ligands for
103
the functionalization and stabilization of inorganic surfaces.
2
To this point, NHCs have
been used as a supporting ligand framework for both colloidal metal nanocrystals ( i.e.,
Au, Pd, Pt, Ru NCs)
3
and planar Au surfaces.
4
As an example of their utility, NHC-
functionalized Au surfaces have been shown to possess both enhanced thermal and
chemical stabilities when compared to thiol functionalization.
4b,5
Additionally, NHC
functionalization has recently been shown to have an important function in increasing the
activity of Au NC electrocatalysts for the reduction of carbon dioxide.
6
Despite the success and demonstrated utility of coupling NHC ligands with
certain metal NCs, such as Au, there has been comparatively less success in using NHCs
to support NCs of other materials.
2,7
For example, while NHC-supported Au NCs are now
well known to possess good colloidal stability,
3a,b
there has heretofore been a noticeable
absence of stable NHC-Cu and NHC-Ag NC congeners in the literature. Likewise, there
has been scant mention in the literature of NHCs being applied to non-metal NCs, such as
covalent semiconductors,
8
even though other L-type ligands such as amines or
phosphines are well known to coordinate to and stabilize many semiconductor NCs.
9

Much of the current lack of material diversity for NHC-stabilized NCs stems from a lack
of fundamental insight into the NHC-NC interface, which has not yet been studied across
different surfaces.
In an effort to better understand the nature of the NHC-NC interface, we herein
developed a synthetic strategy to NHC-stabilized Ag and Ag2E (E = S, Se) NCs in order
to explore NHC ligand binding across both metal and covalent semiconductor NC
surfaces. We observed an empirical difference in colloidal stability between the material
platforms, which was explained in terms of NHC ligand dynamics and binding energies
104
using a complement of solution NMR spectroscopy and density functional theory (DFT)
calculations. This allowed for the extension of this chemistry to the preparation of
colloidal NHC-Cu2– xE NCs, where 0 < x ≤ 1, using NHC-Cu NC intermediates.

4.3. Results and Discussion
4.3.1. Synthesis of NHC-Capped Ag, Ag2S, and Ag2Se NCs
Bromo[1,3-(ditetradecyl)benzimidazol-2-ylidene]silver(I) was used as the NHC
precursor to both NHC-Ag and NHC-Ag2E NCs. This particular NHC has been
previously employed to stabilize colloidal NHC-Au NCs.
3c
Colloidal NHC-Ag NCs were
first obtained via a biphasic reduction of the NHC-AgBr complex with excess NaBH4 in
dichloromethane and water. The resulting NHC-Ag NCs are isolated from the organic
phase, and upon purification, can be redispersed in nonpolar solvents. The lack of stable
NHC-Ag NCs in the literature may be a result of the reduced bond dissociation energy of
the NHC-Ag bonds compared to the NHC-Au bond.
10
As a direct consequence, NHC
ligands may not be as effective at stabilizing Ag NCs since the dissociation of NHC
ligands can result in NC aggregation. Therefore, to achieve colloidally stable NHC-Ag
NCs, a lower concentration of the NHC-AgBr precursor was utilized to prevent
subsequent aggregation of the resulting NCs.
Scheme 4.1 Synthesis of NHC-stabilized colloidal coinage metal and metal chalcogenide NCs.

105

The as-synthesized NHC-Ag NCs can be readily dispersed in non-polar solvents
such as toluene, tetrachloroethylene (TCE), or 1-octadecene (ODE) to give optically
transparent, dark red colloidal suspensions that are stable for approximately one week. A
characteristic localized surface plasmon resonance (LSPR) band from the NHC-Ag NCs
centered at  = 414 nm is observed by UV-vis absorption spectroscopy (Figure 4.1).
Powder X-ray diffraction (XRD) confirms the as-synthesized product crystallizes in the
expected fcc phase of Ag (Figure 4.1b). Transmission electron microscopy (TEM)
analysis reveals the NHC-Ag NCs to be spherical in morphology with a mean diameter of
5.2 nm and standard deviation about the mean (σ/d) of 23% (Figure 4.1c). The size and
size distribution of the NHC-Ag NCs can be readily tuned by changing the reagent
concentrations; for example, if the concentrations of both NHC-AgBr and NaBH4 are
doubled, NHC-Ag NCs with a mean diameter of 9.0 nm (σ/d = 16%) can be achieved.

Figure 4.1 (a) UV-vis absorption spectra of NHC-Ag, Ag 2S, and Ag 2Se NC suspensions in
toluene. Inset is a photograph of dilute toluene suspensions the NCs. (b) Powder XRD patterns of
NHC-Ag, Ag 2S, and Ag 2Se NCs. TEM micrographs of NHC-stabilized (c) Ag, (d) Ag 2S, and (e)
Ag 2Se NCs. Insets are the size distribution of corresponding NCs (n = 500 counts).

Subsequently, the NHC-Ag NCs were used as seeds for the synthesis of NHC-
106
Ag2E NCs (Scheme 4.1). To avoid the introduction of additional extraneous ligands into
the synthesis, the NHC-Ag NCs and S or Se are separately suspended/dissolved in a non-
coordinating solvent (i.e., ODE). NHC-Ag2E NCs can be prepared by injecting the
chalcogen in ODE into the NHC-Ag NC suspension at 120 ˚C. Absorption measurements
display distinctive features into the NIR for the NHC-Ag2S and NHC-Ag2Se NCs, and
confirm the absence of a Ag LSPR band, indicating full conversion from Ag to Ag2E
NCs (Figure 4.1). XRD patterns confirm the purified NC products to be phase-pure
monoclinic Ag2S (PDF no. 00-014-0072) and orthorhombic Ag2Se (PDF no. 00-024-
1041) (Figure 4.1b). The stoichiometry of Ag2S and Ag2Se was confirmed by inductively
coupled plasma-optical emission spectroscopy (ICP-OES), which gives a Ag/S ratio of
2.28 and Ag/Se ratio of 2.11. TEM analysis shows spherical NCs with a mean diameter
of 15.2 nm (σ/d = 17%) for Ag2S and 12.4 nm (σ/d = 17%) for Ag2Se NCs, based on a
5.2 nm NHC-Ag NC seed (Figure 4.1d,e). Additionally, the colloidal stabilities of these
NHC-stabilized Ag2S and Ag2Se NCs are excellent, with no noticeable aggregation or
precipitation after >6 months.

4.3.2. Characterization of Surface NHC Ligands
To understand the marked differences in colloidal stability between the NHC-
stabilized metal and covalent semiconductor surfaces, the colloidal NHC-Ag and Ag2S
NCs were interrogated as model systems using a battery of spectroscopic techniques.
Solution
1
H NMR was applied to gain direct insight into the chemical identity of the
surface ligands. The NHC-Ag NCs display a similar
1
H NMR spectrum to the NHC-
AgBr precursor (Figure 4.2a), with a single set of benzimidazole and tetradecyl
107
resonances. However, the observed broadening of these resonances for the NHC-Ag NCs
results from slower tumbling of the ligands, indicating that the NHC is coordinated to the
NC surface.
11
The same set of resonances was also observed for the NHC-Ag2S NCs,
with a slight downfield shift of all corresponding peaks (Figure 4.2a). Further proof
follows from a direct comparison of the
1
H 
13
C HSQC spectra of both NHC-AgBr and an
NHC-Ag NC suspension (Figure 4.2b,c). The NHC-Ag NC HSQC cross peaks indeed
agree well with those of the corresponding NHC-AgBr complex. These spectral data
collectively suggest that the NHC ligands are coordinating to the surfaces of both the Ag
and Ag2S NCs.

Figure 4.2 (a)
1
H NMR spectrum of NHC-AgBr, NHC-Ag NCs, and NHC-Ag 2S NCs.
Resonances from 0.2 to 8 ppm are assigned accordingly. Solvent impurities are indicated by
*(CH 2Cl 2), (toluene), (ethanol), and •(H 2O). (b,c)
1
H 
13
C HSQC spectra of NHC-AgBr and a
colloidal Ag NCs suspension in CDCl 3, respectively. (d f) High-resolution XPS spectra of N 1s
for NHC-Ag, Ag 2S, and Ag 2Se NCs, respectively.

To gain a more comprehensive picture of the NHC-NC interface, diffusion
ordered NMR spectroscopy (DOSY) was performed on suspensions of both purified
NHC-Ag and NHC-Ag2S NCs in CDCl3 (Figure 4.3). Diffusion coefficients for the NHC
108
ligands of 5.0  10
–10
m
2
/s and 2.3  10
–10
m
2
/s were measured for Ag NCs and Ag2S
NCs, respectively. Both of these diffusion coefficients are lower than the free NHC-AgBr
complex, which exhibits a diffusion coefficient of 6.5  10
–10
m
2
/s. The hydrodynamic
diameter of the NCs can be obtained from the diffusion coefficients (d H) based on the
Stokes-Einstein equation (dH =
𝑘 𝐵 𝑇 3𝜋 𝐷 (kB, Boltzmann constant; T, absolute temperature; ,
solvent viscosity, 0.57 mPa s for CDCl3);
11
however, the observed dH of the NHC-Ag and
NHC-Ag2S NCs were only 1.7 and 3.7 nm, respectively, which are smaller than the
expected core diameters of 5 and 15 nm based on TEM. This thus suggests a dynamic
NHC-NC interface, with an equilibrium between bound and free NHC ligands (likely
NHC-Ag complexes). As a result, a weighted mean diffusion coefficient between the
bound and unbound NHC ligands was observed.
12
On this basis, NHC ligands on the
surface of both Ag and Ag2S NCs appear to be labile; however, this does not explain the
observed difference in colloidal stability between the Ag and Ag2S NCs.

Figure 4.3 DOSY spectra of (a) NHC-AgBr, (b) Ag NCs, and (c) Ag 2S NCs in CDCl 3 solution.
Diffusion coefficients were measured to be 6.5  10
–10
, 5.0  10
–10
, and 2.3  10
–10
m
2
/s for NHC-
AgBr, Ag NCs and Ag 2S NCs, respectively.

The surface chemistry of the NHC-NCs was further corroborated by X-ray
photoelectron spectroscopy (XPS), which shows the presence of N on the surface of the
NHC-Ag, Ag2S, and Ag2Se NCs through a N 1s peak at a binding energy of ~400.5 eV
(Figure 4.2d-f). Interestingly, Br was revealed to be present in only the NHC-Ag NCs (as
109
had been previously observed for NHC-Au NCs
3a
), but not in the case of the NHC-Ag2S
or Ag2Se NCs (Figure 4.4). The atomic Br/Ag ratio is quantified by XPS to be <14% for
the NHC-Ag NCs, thus the surface is not expected to be entirely shelled with Br. This
degree of surface Br does, however, suggest that Br plays a role at the surface of the Ag
NCs, while only covalent NHC-Ag(I) interactions are present in the NHC-Ag2E NCs.
Moreover, the presence of Br in the NHC-Ag NCs is not associated with protonated
benzimidazolium bromide since a proton resonance at ~11.6 ppm is not observed in the
1
H NMR spectrum of the NCs.

Figure 4.4 High-resolution XPS spectra of (a) NHC-Ag, (b) NHC-Ag 2S, and (c) NHC-Ag 2Se
NCs, respectively. Note the absence of Br 3d signals from NHC-Ag 2S and NHC-Ag 2Se NCs,
indicating the covalent interaction between the NHC ligands and Ag 2E NC surface.

The presence of NHC ligands was further confirmed by FT-IR spectra, which
show characteristic bands for imidazolium ring stretching at 1467 cm
1
(s, sym) in NHC-
110
AgBr, NHC-Ag, NHC-Ag2S, and NHC-Ag2Se NCs (Figure 4.5a). The quantity of NHC
on the NC surfaces was determined by thermogravimetric analysis (Figure 4.5b), with the
mass loss at ~320 ˚C assigned to the loss of NHC ligands.
13
The NHC-Ag and NHC-
Ag2E NCs display a single-step mass loss event at ~320 ˚C of 40%, 3.4%, and 8.6% for
NHC-Ag, Ag2S, and Ag2Se NCs, respectively. The number of NHC ligands/NC can then
be estimated assuming a spherical morphology for Ag (5 nm), Ag2S (15 nm), and Ag2Se
(12 nm) NCs to be 603, 530 and 825, respectively. These values are all below the
reported surface density previously reported for 4-nm NHC-Au NCs (~730 ligands per
NC).
13

Figure 4.5 (a) FT-IR and (b) TGA spectra of NHC-AgBr, NHC-Ag, NHC-Ag 2S, and NHC-
Ag 2Se NCs.

4.3.3. Energetic NHC-NC Interface Exposed from Density Functional Theory (DFT)
Calculation
Since the dynamic nature of the NHC-NC interface does not appear to be the
origin of the difference in colloidal stability between the Ag and Ag2E NCs, we
111
investigated potential differences in binding energy. The binding energies of a model 1,3-
(dimethyl)benzimidazol-2-ylidene NHC ligand to ~1.2 nm clusters of Ag and Ag2S were
calculated using DFT. Because of the experimentally observed presence of Br in the Ag
NCs, we modified the Ag cluster surface with Br atoms. The calculations revealed that
the binding interaction between the NHC and the Br-modified Ag cluster is weak. The
NHC molecule could only bind to the vertex site on the surface of the Ag cluster with a
small binding energy of 0.30 eV (Figure 4.6a). The vertex site of the Ag NC is least
coordinated, providing the best opportunity for ligand binding. For the other two binding
sites, namely the fc and ridge sites, the NHC failed to bind to the surface of the Ag cluster
with negligible binding energies and large distances between the coordinating C atom in
NHC and the Ag atom on the cluster (Figure 4.6b,c). The reason may be attributed to the
steric hindrance from the alkyl groups in the NHC interfering with the surface Br. The
calculated Bader charge in the vertex site shows that the electron charge was transferred
from surface Ag to Br, giving about 0.43 e/Br atom. Accordingly, a weak modification of
the C orbitals is observed in the valence band of the density of states (DOS) (Figure
4.7a). In comparison, the NHC ligand binds to the Ag2S cluster through relatively strong
bonding interactions between the carbene C and Ag
+
cations on the surface of the cluster.
The geometries and binding energies for three types of binding of the NHC to the Ag2S
cluster are similar (Figure 4.6d-f), suggesting a similar local environment for the Ag
atoms on the surface. The binding energy varies from 0.57 to 0.77 eV for the three
binding geometries with the C–Ag bond decreasing from 2.47 to 2.26 Å, respectively.
The relatively strong binding of the NHC to the Ag2S cluster is also reflected in the DOS
for the geometry of binding site-2, where a strong modification of the C orbitals is
112
observed in the valence band (Figure 4.7a).

Figure 4.6 Geometries and binding energies for three types of binding of an NHC ligand to a (a-
c) Ag cluster that was decorated with four Br atoms surrounding the binding site, and a (d-f) Ag 2S
cluster. Ag, S, Br, H, C and N atoms were depicted in grey, yellow, brown, pink, deep and light
blue, respectively. The red Ag atom denotes the surface-binding site for the NHC ligand.

It is noteworthy that surface passivation by NHC ligands does not appear to
introduce any mid-gap trap states in the Ag2S cluster since the orbitals of the NHC
ligands only appear in the valence or conduction bands. Therefore, as a strong donor, the
neutral NHC ligands not only provide a strong binding energy to the Ag2S NCs to impart
excellent colloidal stability, but also present the possibility to positively affect the
optoelectronic properties of covalent semiconductor NCs.
113

Figure 4.7 (a) Density of states for the vertex binding geometry of the NHC molecule binding to
the Ag cluster decorated with Br atoms (upper panel) and the isolated NHC molecule (lower
panel). The dashed line denotes the Fermi level. The most inner orbitals of the NHC molecule are
aligned between the two systems because they are only mildly affected when the NHC molecule
binds to the Ag cluster. The binding occurs through the C atom of the ligand. The C orbitals
remain nearly intact in the valence band (rectangular) region, because binding of the NHC
molecule to the Ag cluster is weak. (b) Spin-polarized density of states for the site-2 binding
geometry of the NHC molecule binding to the Ag 2S cluster (upper panel) and the isolated NHC
molecule (lower panel). The dashed line denotes the Fermi level. The C orbitals of the ligand are
significantly modified in the valence band (rectangular) region due to strong NHC-Ag 2S
interaction.

4.3.4. Synthesis of NHC-Capped Cu, Cu2 xS, and Cu2 xSe NCs
Furthermore, the same chemistry was extended to NHC-CuBr to generate well-
defined NHC-Cu2–xE NCs (0 < x ≤ 1, Figure 4.8). The reduced product from NHC-CuBr
is shown to be fcc Cu NCs, which possess limited oxidative and colloidal stability.
Subsequent DFT calculations further show a very weak binding interaction between the
NHC ligand and an analogous Cu cluster decorated with Br (Figure 4.9). After dispersing
the NHC-Cu NCs in ODE, colloidal NHC-Cu2–xE NCs can be prepared following the
same reaction procedure as before. Interestingly, varying the nominal chalcogen
stoichiometry in the reaction allows for control over the copper chalcogenide phase. For
instance, hexagonal NHC-CuE NCs were produced if chalcogen was added in a 1:1
114
(Cu/E) molar ratio, while cubic NHC-Cu1.8E NCs were obtained with a 2:1 (Cu/E) molar
ratio. The addition of this NHC-MBr precursor chemistry to give NHC-Cu and Cu2–xE
NCs further extends the generality of this approach in generating new metal and covalent
semiconductor NCs that are coordinated by NHC ligands. As with their Ag counterparts,
the NHC-Cu2–xE NCs possess superior colloidal stability than the NHC-Cu NC
intermediates.

Figure 4.8 (a) Powder XRD patterns of NHC-Cu, Cu 2–xS, and Cu 2–xSe NCs. The asterisk in the
diffraction pattern for the Cu NCs indicates the presence of trace Cu 2O. (b) UV-vis-NIR
absorption spectra of NHC-Cu, Cu 2–xS, and Cu 2–xSe NC suspensions in TCE. Inset is a
photograph of dilute suspensions of Cu, CuS, CuSe, Cu 1.8S and Cu 1.8Se NCs in TCE. (c-f) TEM
115
micrographs of NHC-CuS (d = 11.7  1.8 nm), CuSe (d = 9.4  1.3 nm), Cu 1.8S (d = 7.4  0.8
nm) and Cu 1.8Se (d = 9.5  1.1 nm) NCs, respectively.

Figure 4.9 Geometry and binding energy for the vertex binding of an NHC ligand to the Cu
cluster which was decorated with four Br atoms. Cu, Br, H, C and N atoms were depicted in
orange, brown, pink, deep and light blue. The red atom denotes the Cu atom at binding site.

4.4. Experimental
4.4.1. General Information and Characterization Techniques
Reagents and solvents were purchased from commercial sources and used as received,
unless otherwise stated. Benzimidazole (C7H6N2, 99%), 1-bromotetradecane (C14H29Br),
silver( ) oxide (99+%), copper( ) oxide (99.9%), potassium carbonate (K2CO3, anhydrous,
99%), 1,4-dioxane, and selenium powder (-200 mesh, 99.999%) were purchased from
Alfa Aesar. Sodium borohydride (NaBH4, 99%), 1-octadecene (ODE, technical grade,
90%) and sulfur powder (99.98%) were purchased from Sigma-Aldrich. Reactions
involving air- or moisture-sensitive compounds were conducted under a nitrogen
atmosphere by using standard Schlenk techniques.
UV-vis-NIR spectroscopy was carried out on a Perkin-Elmer Lamba 950
spectrophotometer equipped with a 150 mm integrating sphere, using a quartz cuvette for
liquid samples.
116
Thermogravimetric analysis (TGA) measurements were made on a TA
Instruments TGA Q50 instrument, using sample sizes between 5-10 mg in an alumina
crucible under a flowing nitrogen atmosphere. TGA samples were prepared by drying
the colloid under flowing nitrogen at 80 ˚C for up to 120 min, then lightly crushing with a
spatula prior to analysis.
FT-IR spectra were acquired from pressed pellets on a Bruker Vertex 80. Pressed
pellets were made of dried nanocrystals (~2 mg) in a dry KBr matrix (~100 mg).
Solution 1D and 2D
1
H,
13
C,
1
H 
13
C HSQC NMR spectra were collected at
ambient temperature on a Varian 500 spectrometer (500 MHz in
1
H) with chemical shifts
represented in units of ppm. All spectra were referenced to the residual solvent peaks
(5.33 ppm for CD2Cl2, 7.26 ppm for CDCl3). NMR samples were prepared by drying the
nanocrystal colloid under nitrogen flow and washing by ethanol 3  before dispersing in
CD2Cl2 or CDCl3. For molecular complexes, 32 scans with a delay time of 1 s were
taken for each sample, and the data are presented as averages of those scans. For
nanocrystal samples, 1024 scans with a delay time of 30 s were taken to allow complete
relaxation between pulses for
1
H spectra. Typical NMR samples had a nanocrystal
concentration in the range of 10 30 M. All diffusion measurements were made using
the pulse sequence Dbppste_cc (Bipolar Pulse Pair STimulated Echo with Convection
Compensation). The strength of the gradient was calibrated by measuring the self-
diffusion of the residual HDO signal in a D2O sample at 298 K (1.90  10
–9
m
2
s
–1
).
14
A
diffusion delay  of 200 ms and diffusion gradient length  of 5 ms were used for
nanocrystal samples. The amplitude of the gradient pulses was varied from 2% to 90% of
the maximum amplitude of 16.7 G cm
–1
in 32 steps. Diffusion coefficients can be
117
subsequently obtained from Vnmrj’s analysis routines. Hydrodynamic diameter d H was
then calculated using the Stokes-Einstein equation: dH =
𝑘 𝐵 𝑇 3𝜋 𝐷 (kB, Boltzmann constant; T,
absolute temperature; , solvent viscosity, 0.57 mPa s for CDCl3).
11

X-ray photoelectron spectra (XPS) were obtained using a Kratos Axis Ultra X-ray
photoelectron spectrometer with an analyzer lens in hybrid mode. High resolution scans
were performed using a monochromatic aluminum anode with an operating current of 6
mA and voltage of 10 kV using a step size of 0.1 eV, a pass energy of 40 eV, and a
pressure range between 1 3  10
–8
torr. The binding energies for all spectra were
referenced to the C1s core level at 284.8 eV.
Inductively coupled plasma-optical emission spectroscopy (ICP-OES) was
performed on the nanocrystal samples using a Thermo Scientific Icap 7000 series ICP-
OES. Nanocrystal samples were prepared by drying out 10 mg of solid, digesting in
concentrated HNO3, and diluting with ca. 99 mL 10% HNO3. Silver to chalcogen (S or
Se) ratio was calculated by averaging 3 independently synthesized nanocrystal samples.
Elemental analysis was conducted by a Flash 2000 Elemental Analyzer (CHNS
Elemental Analyzer). Samples (2 3 mg) were prepared with tin crucibles.
Powder X-ray diffraction (XRD) data was collected using a Rigaku Ultima IV
diffractometer in parallel beam geometry (2 mm beam width) using Cu K  radiation (  =
1.54 Å). Samples were prepared by drop casting onto zero-diffraction, single crystal Si
substrates, followed by moderate heating to evaporate solvents.
Transmission electron microscopy (TEM) analysis was performed on a JEOL
JEM-2100 microscope at an operating voltage of 200 kV, equipped with a Gatan Orius
118
CCD camera. Samples for TEM analysis were prepared from dilute purified nanocrystal
samples deposited on 400 mesh carbon-coated copper grids (Ted Pella, Inc.).

4.4.2. Synthesis of Benzimidazolium Salts, NHC-AgBr, and NHC-CuBr Complexes
Scheme 4.2 General Synthetic Scheme

Benzimidazolium salt (1), NHC-AgBr (2), and NHC-CuBr (3) complexes were
synthesized according to modified literature procedures.
15
The following abbreviations
are used: s = singlet, d = doublet, t = triplet, q = quadruplet, m = multiplet, br = broad.
119
1,3-(Ditetradecyl)benzimidazolium bromide (1).
15a
Benzimidazole (2.36 g, 20 mmol),
K2CO3 (2.76 g, 20 mmol), n-tetradecyl bromide (18 mL, 60
mmol) and CH3CN (20 mL) were added into a three-neck flask
and stirred at reflux (~85 C) for 24 h. After the reaction, the solvent was removed under
reduced pressure, followed by the dissolution in CH2Cl2. The mixture was filtered to
remove the precipitate (KBr). The filtrate was then concentrated under reduced pressure.
After the residue was recrystallized from CH2Cl2/pentane and dried under vacuum, a
white solid (6.8 g, 57%) was obtained.
1
H NMR (400 MHz, CDCl3)  11.56 (s, 1H),
7.71 7.63 (m, 4 H), 4.62 (t, J = 7.6 Hz, 4H), 2.10 2.00 (m, 4H), 1.45 1.2 (m, 44H), 0.87
(t, J = 7.1 Hz, 6H).
13
C NMR (101 MHz, CDCl3)  143.06, 131.46, 127.18, 113.18,
47.82, 32.04, 29.79, 29.76, 29.70, 29.68, 29.62, 29.51, 29.47, 29.17, 22.81, 14.24 
 C-AgBr (2).
15a
To a solution of 1 (1.2 g, 2.0 mmol) in dried CH2Cl2 (40 mL), Ag2O
(0.56 g, 2.4 mmol) was added. The mixture was refluxed for 20
h, and excess Ag2O was filtered away. The filtrate was
concentrated under reduced pressure. After the residue was
recrystallized from CH2Cl2/pentane and dried under vacuum, a light brown solid (0.94 g,
74%) was obtained.
1
H NMR (500 MHz, CD2Cl2)  7.51 (dd, J = 5.9 and 3.0 Hz, 2H),
7.43 (dd, J = 6.0 and 3.2 Hz, 2 H), 4.62 (t, J = 7.3 Hz, 4H), 2.101.9 (m, 4H), 1.4 1.1 (m,
44H), 0.87 (t, J = 6.7 Hz, 6H).
13
C NMR (125 MHz, CD2Cl2)  134.27, 124.41, 112.14,
50.14, 32.50, 30.81, 30.26, 30.23, 30.20, 30.12, 30.04, 29.93, 29.81, 27.40, 23.27, 14.45.
Elemental analysis: measured % (theoretical %) for C35H62N2AgBr: C: 60.56 (60.17), N
4.10 (4.01), H 9.18 (8.95).
120
NHC-CuBr (3).
15b
Cu2O (0.34 g, 2.4 mmol) was added to a solution of 1 (1.2 g, 2.0
mmol) in dried 1,4-dioxane (40 mL). The mixture was refluxed
for 20 h, and excess Cu2O was filtered. The filtrate was
concentrated under reduced pressure. The residue can be isolated
by filtration, and washed by hexanes. After drying under vacuum, the title compound
was obtained as a brown solid (0.9 g, 74%).
1
H NMR (500 MHz, CD2Cl2)  7.48 (dd, J =
6.0 and 3.0 Hz, 2H), 7.40 (dd, J = 6.0 and 3.0 Hz, 2 H), 4.62 (t, J = 7.2 Hz, 4H), 2.10 1.9
(m, 4H), 1.4 1.1 (m, 44H), 0.87 (t, J = 6.8 Hz, 6H).
13
C NMR (125 MHz, CD2Cl2)
 124.25, 111.94, 49.49, 32.50, 30.82, 30.26, 30.23, 30.20, 30.12, 30.04, 29.93, 29.78,
27.38, 23.27, 14.45. Elemental analysis: measured % (theoretical %) for
C35H62N2CuBr•(H2O)2: C: 60.45 (60.91), N 4.58 (4.06), H 9.26 (9.64).

4.4.3. Synthesis of NHC-Metal and NHC-Metal Chalcogenide NCs
NHC-Ag NCs. 2 (140 mg, 0.2 mmol) was first dissolved in 80 mL CH2Cl2 at 0.4 ˚C (ice
bath) in air, giving a 2.5 mM solution. To the solution of 2, a solution of NaBH4 (48 mg,
1.2 mmol) in DI-H2O (80 mL) was added dropwise while maintaining the ice temperature.
The mixture changed color from colorless to orange and red upon borohydride addition,
indicating the formation of Ag NCs. After dropwise addition (~20 min), the mixture was
allowed to stir at room temperature for 100 min. The CH2Cl2 layer was then separated
from the aqueous layer and washed by DI-H2O twice before the solvent was removed
under reduced pressure. The residue was washed by ethanol and the crude suspension
was centrifuged. The supernatant (EtOH) was removed. After drying under flowing
121
nitrogen, the as-synthesized NHC-Ag NCs as a dark powder were used for XRD, XPS,
TGA, FT-IR, and NMR analysis. NHC-Ag NCs were also redispersed in toluene to give
an optically transparent,s dark red colloidal suspension that was used for UV-vis-NIR
absorption and TEM analysis.
NHC-Cu NCs. Since the reduction potential of Cu
+
/Cu (E =0.52 V) is lower than Ag
+
/Ag
(E =0.80 V), the reduction of 3 by NaBH4 was conducted at room temperature instead of
0 ˚C, as for the Ag NCs. 3 (130 mg, 0.2 mmol) was first dissolved in 40 mL CH2Cl2 at
room temperature under a nitrogen atmosphere to give a 5 mM solution. A nitrogen
atmosphere was used here to prevent the resulting NHC-Cu NCs from oxidation. To the
solution of 3, a solution of NaBH4 (48 mg, 1.2 mmol) in DI-H2O (40 mL) was added
dropwise at room temperature. After addition (~10 min), the mixture was allowed to stir
at room temperature for another 10 min. The CH2Cl2 layer was then separated from the
aqueous layer and washed by DI-H2O twice before the solvent was removed by reduced
pressure. The residue was washed by ethanol and the crude suspension was centrifuged.
The supernatant (EtOH) was removed. After drying under nitrogen flow, the as-
synthesized NHC-Cu NCs as a dark powder were used for XRD, XPS, TGA, FT-IR, and
NMR analysis. NHC-Cu NCs were also redispersed in toluene to give an optically
transparent, light brown colloidal suspension that was used for UV-vis-NIR absorption
and TEM analysis.
NHC-M2 –xE NCs (M = Ag, Cu; E = S, Se). The as-prepared dried NHC-M NCs were first
dispersed in ODE (10 mL) at room temperature under a nitrogen atmosphere. The
temperature was then raised to 120 ˚C. Elemental S or Se powder was suspended in ODE
under a nitrogen atmosphere to give a 0.1 M chalcogen suspension. The chalcogen
122
suspension with a 2:1 (M:E) or 1:1 (M:E) stoichiometry was rapidly injected to NHC-M
NCs suspension at 120 ˚C. After 10 min, hexanes (10 mL) were injected to thermally
quench the reaction. The reaction was then cooled to room temperature, followed by
centrifugation to remove excess Se or aggregated NCs. NHC-M2–xE NCs were then
isolated by precipitation from excess acetone and redispersion in toluene or TCE.

4.5. Conclusions
We have successfully demonstrated the synthesis of colloidally stable NHC-stabilized Ag
metal and Ag2E metal chalcogenide NCs using a bromo[1,3-(ditetradecyl)benzimidazol-
2-ylidene]silver(I) synthon. With a combination of
1
H NMR, HSQC, and DOSY spectra,
a dynamic coordination sphere of effective but labile NHC ligands is found to coordinate
to both Ag and Ag2S NCs, with the Ag2S NCs observed to be more colloidally stable
with respect to time. This empirical difference is rationalized to be a result of the
different binding energies between the NHC ligand and NC surface for the two materials,
with a significantly larger binding affinity between the NHC and Ag2S NCs revealed by
DFT calculations. These experimental and computational results help elucidate a more
comprehensive understanding of the NHC-NC interface, providing new insight into the
surface chemistry of colloidal NCs with L-type NHC ligands across a range of materials.
The excellent colloidal stability of NHC-M2–xE NCs, and implications of strong donation
with respect to their optoelectronic properties, suggests these are promising ligands to be
further explored for the stabilization of covalent semiconductor NCs.

123
4.6. References

(1) Hopkinson, M. N.; Richter, C.; Schedler, M.; Glorius, F. Nature 2014, 510, 485-
496.
(2) Zhukhovitskiy, A. V.; Mavros, M. G.; Van Voorhis, T.; Johnson, J. A. J. Am.
Chem. Soc. 2013, 135, 7418-7421.
(3) (a) Vignolle, J.; Tilley, T. D. Chem. Commun. 2009, 7230-7232; (b) Ranganath,
K. V. S.; Kloesges, J.; Schä fer, A. H.; Glorius, F. Angew. Chem. Int. Ed. 2010, 49,
7786-7789; (b) Ling, X.; Roland, S.; Pileni, M.-P. Chem. Mater. 2015, 27, 414-
423; (c) Ranganath, K. V. S.; Kloesges, J.; Schä fer, A. H.; Glorius, F. Angew.
Chem. Int. Ed. 2010, 49, 7786-7789; (d) Lara, P.; Rivada-Wheelaghan, O.;
Conejero, S.; Poteau, R.; Philippot, K.; Chaudret, B. Angew. Chem. Int. Ed. 2011,
50, 12080-12084; (e) Baquero, E. A.; Tricard, S.; Flores, J. C.; de Jesú s; E.;
Chaudret, B. Angew. Chem. Int. Ed. 2014, 53, 13220-13224.
(4) (a) Zhukhovitskiy, A. V.; Mavros, M. G.; Van Voorhis, T.; Johnson, J. A. J. Am.
Chem. Soc. 2013, 135, 7418-7421; (b) Crudden, C. M.; Horton, J. H.; Ebralidze, I.
I.; Zenkina, O. V.; McLean, A. B.; Drevniok, B.; She, Z.; Kraatz, H.-B.; Mosey,
N. J.; Seki, T.; Keske, E. C.; Leake, J. D.; Rousina-Webb, A.; Wu, G. Nat. Chem.
2014, 6, 409-414.

(5) Ling, X.; Schaeffer, N.; Roland, S.; Pileni, M.-P. Langmuir 2015, 31, 12873-
12882.
(6) Cao, Z.; Kim, D.; Hong, D.; Yu, Y.; Xu, J.; Lin, S.; Wen, X.; Nichols, E. M.;
Jeong, K.; Reimer, J. A.; Yang, P.; Chang, C. J. J. Am. Chem. Soc. 2016, 138,
8120-8125.

(7) Ling, X.; Schaeffer, N.; Roland, S.; Pileni, M.-P. Langmuir 2013, 29, 12647-
12656.
(8) Aslanov, L.; Zakharov, V.; Zakharov, M.; Kamyshny, A.; Magdassi, S.;
Yatsenko, A. Russ. J. Coord. Chem. 2010, 36, 330-332.
(9) Brutchey, R.; Hens, Z.; Kovalenko, M. V. Chemistry of Organo-Hybrids:
Synthesis and Characterization of Functional Nano-Objects, eds. B. Charleux, C.
Copé ret, E. Lacô te, John Wiley & Sons, Inc., Hoboken, NJ, 2015, ch. 7.
(10) Jacobsen, H.; Correa, A.; Poater, A.; Costabile, C.; Cavallo, L. Coord. Chem. Rev.
2009, 253, 687-703.
(11) Hens, Z.; Martins, J. C. Chem. Mater. 2013, 25, 1211-1221.
124
(12) (a) Grisorio, R.; Debellis, D.; Suranna, G. P.; Gigli, G.; Giansante, C. Angew.
Chem. Int. Ed. 2016, 55, 6628-6633; (b) De Roo, J.; Ibá ñ ez, M.; Geiregat, P.;
Nedelcu, G.; Walravens, W.; Maes, J.; Martins, J. C.; Van Driessche, I.;
Kovalenko, M. V.; Hens, Z. ACS Nano 2016, 10, 2071-2081.
(13) MacLeod, M. J.; Johnson, J. A. J. Am. Chem. Soc. 2015, 137, 7974-7977.
(14) Cros-Gagneux, A.; Delpech, F.; Nayral, C.; Cornejo, A.; Coppel, Y.; Chaudret, B.
J. Am. Chem. Soc. 2010, 132, 18147-181577.
(15) (a) Ling, X.; Roland, S.; Pileni, M.-P. Chem. Mater. 2015, 27, 414-42; (b) Chun,
J.; Lee, H. S.; Jung, I. G.; Lee, S. W.; Kim, H. J.; Son, S. U. Organometallics
2010, 29, 1518-1521.

125
Chapter 5. Tunable Room-Temperature Synthesis of Coinage Metal
Chalcogenide Nanocrystals from N-Heterocyclic Carbene Synthons*
*Published in Chem. Mater. 2017, 29, 1396–1403.

5.1. Abstract
We present a new toolset of precursors for semiconductor nanocrystal synthesis, N-
heterocyclic carbene (NHC)-metal halide complexes, which enable a tunable molecular
platform for the preparation of coinage metal chalcogenide quantum dots (QDs). Phase-
pure and highly monodisperse coinage metal chalcogenide (Ag2E, Cu2–xE; E = S, Se)
QDs are readily synthesized from the direct reaction of an NHC-MBr synthon (where M
= Ag, Cu) with alkylsilyl chalcogenide reagents at room temperature. We demonstrate
that the size of the resulting QDs is well tailored by the electron-donating ability of the L-
type NHC ligands, which are further confirmed to be the only capping ligands on the QD
surface, imparting excellent colloidal stability. Local superstructures of the NHC-capped
Ag2S QDs are observed by TEM, further demonstrating their potential for synthesizing
monodisperse ensembles and mediating self-assembly.

5.2. Introduction
Semiconductor nanocrystals, or quantum dots (QDs), are attractive functional materials
for photovoltaics, photocatalysis, displays, and biomedical applications because of their
uniquely tunable optoelectronic properties.
1-5
In particular, the coinage metal
chalcogenide (i.e., Ag2E, Cu2–xE; E = S, Se) QDs have gained significant attention
resulting from their band gaps (Eg = 0.15 2.0 eV), low toxicity (as compared to Cd- and
126
Pb-containing semiconductors), high absorption coefficients, NIR surface plasmon
oscillation, and NIR emission.
6,7
Coinage metal chalcogenide QDs have proven to be
useful for photovoltaics,
8
bioimaging,
9
thermochromics,
10
NIR absorption,
11
and
thermoelectrics.
12,13
Moreover, these QDs also have served as parent platforms for other
chalcogenide-based ternary (CuInS2, Cu2SnS3)
14
and quaternary (Cu2ZnSnS4)
15

semiconductor nanocrystals via cation exchange or as reaction intermediates.
Conventional synthetic routes to these coinage metal chalcogenide QDs often involve
inorganic salts (e.g., AgBr, CuCl, AgNO3) that are solvated at elevated temperatures by
long-chain organic ligands (e.g., fatty carboxylic acids, amines, phosphines, thiols).
9,16-19

Nanocrystal size and morphology is often controlled by the divergent stages of nucleation
and growth through the LaMer mechanism in so-called hot-injection syntheses.
20,21
While
these high-temperature synthetic approaches have been extensively explored for many
years, there are only a limited number of reports on the room-temperature synthesis of
coinage metal chalcogenide QDs.
22-24
Preparation of functional QDs under room-
temperature conditions is extremely tempting from the perspective of cost and energy
efficiency;
22,25
however, it still remains a challenge to control and tune QD size and size
distribution under ambient conditions, as compared to high-temperature syntheses.
Prasad’s group has recently reported a room temperature synthesis of covellite
phase CuS nanocrystals using CuCl2 dissolved in oleylamine reacted via multiple
injections of (NH4)2S.
22
Well-defined CuS nanocrystals were synthesized by very
carefully tailoring the number of (NH4)2S injections, the amount of (NH4)2S per injection,
and the time interval between each injection. On the other hand, the preparation of phase-
pure digenite Cu1.8S nanocrystals with high monodispersity under ambient conditions
127
remains unreported, and more importantly, there is no general synthetic approach that
gives monodisperse coinage metal chalcogenide (i.e., Ag2E, Cu2–xE; E = S, Se)
nanocrystals at room temperature. At the origin of this challenge is the lack of precursors
with appropriate solubility and reactivity that enable the room-temperature synthesis of
colloidal coinage metal chalcogenide QDs with well-defined composition and
morphology.
In the search for synthetic conditions that might afford colloidal coinage metal
chalcogenide QDs under ambient conditions, we were inspired by the recent examples of
metal nanocrystal syntheses from N-heterocyclic carbene-metal halide complexes (NHC-
MXn),
26-32
as these might serve as precursors with potentially tunable reactivity for QD
synthesis. In this article, we present a new synthetic strategy to coinage metal
chalcogenide QDs under ambient conditions with various types of NHC ligands. NHCs
are a useful and versatile class of L-type (neutral) σ-donating ligands, and a wide
assortment of NHC-metal complexes has been explored in organometallic chemistry.
33-36

With the chemical stability and wide structural tunability of NHC ligands, NHC-metal
complexes may bring new opportunities to colloidal nanocrystal synthesis. For instance,
by tuning the steric and/or electronic properties of NHC ligands, one should, in theory, be
able to tailor the kinetic parameters of synthesis,
37
leading to various particle sizes and/or
shapes, as well as tunable surface functionality assuming the NHC is the only viable
ligand in the synthesis.
32
Thus far, NHC-metal complexes have only been utilized in the
direct preparation of metal nanocrystals (i.e., Au, Pd, Pt, Ru).
26-32
The size,
26

polydispersity,
30
and catalytic activity
38
tunability of the resulting NHC-metal nanocrystal
constructs has been demonstrated with a variety of NHC ligands. These NHC-metal
128
synthons have only recently been explored for the synthesis of QDs; however, this was
only achieved through the chalcogenization of an NHC-metal nanocrystal intermediate.
39

NMR spectroscopy and DFT calculations revealed that the NHC ligands are more
effective ligands in terms of imparting colloidal stability for metal chalcogenide
nanocrystals as compared to their metal nanocrystal counterparts. Herein, we develop a
one-step, room temperature synthesis of monodisperse Ag2E and Cu2–xE QDs from NHC-
MBr precursors (M = Ag, Cu) with excellent reproducibility. The QD size dependence
with respect to electronic and steric parameters of the NHC ligands is investigated, giving
kinetic control over the QD synthesis from the molecular precursor rather than time or
temperature. This route to coinage metal chalcogenide QDs opens the door to
semiconductor nanocrystal synthesis from a wide array of organometallic NHC
complexes.

5.3. Results and Discussion
5.3.1. Synthesis of NHC-Capped Ag2S QDs
It has previously been shown that the reaction of alkylsilyl chalcogenides ((RSi) 2E; E =
S, Se) with metal halides (MXn) in solution yields bulk metal chalcogenides under
ambient conditions.
41
This metathesis reaction can be explained by the hard and soft
(Lewis) acid and base (HSAB) principle (Scheme 5.1), with alkylsilyl chalcogenides
therefore being highly reactive toward transition metal halides. Here, we arrest growth in
this reaction to yield coinage metal chalcogenide QDs by the introduction of NHC
ligands with long-chain n-alkyl groups at the N-substituents.

129
Scheme 5.1 Synthesis of (a) NHC-AgBr complex and (b) NHC-Ag 2S QDs.

The starting NHC-AgBr complex can be readily prepared from the reaction of the
NHC-bromide salts and Ag2O (Scheme 5.1a). The NHC-AgBr complex bromo[1,3-
(ditetradecyl)benzimidazol-2-ylidene] silver( ) (1-b-AgBr) with N-substituents of C14H29
was first employed to investigate the synthesis of Ag2S QDs under ambient conditions.
Upon reaction of 1-b-AgBr with (TMS)2S in dichloromethane, a dark red colloidal
suspension formed immediately at room temperature, indicating a fast nucleation process.
We propose that the reaction proceeds through NHC-stabilized silver sulfido complexes
(such as NHC-AgSAg-NHC, Scheme 5.1b), as similar silver and copper
phenylchalcogenolate complexes have been previously observed,
42
followed by the
nucleation of the monomer species. Upon completion, the resulting QDs were confirmed
by powder X-ray diffraction (XRD) to be phase-pure monoclinic Ag2S (PDF no. 00-014-
0072, Figure 5.1a). Reaction progress was monitored by UV-vis-NIR absorption and
TEM analysis (Figure 5.2). The UV-vis-NIR spectra display featureless absorption into
130
the NIR for the Ag2S QDs, with no observable absorption difference after 5 min. TEM
analysis shows highly monodisperse, spherical QDs, with a gradual growth and focusing
of particle size over the reaction from 7.8  0.8 nm (5 min), to 8.2  0.7 nm (30 min), to
10.3  0.6 nm (60 min) (Figure 5.2). These observations suggest fast nucleation with a
slower stage of nanocrystal growth. The as-prepared Ag2S QDs possess an excellent size
distribution (standard deviation/mean diameter, /d = 6%) after a 60 min reaction (Figure
5.1d f). High-resolution TEM (HRTEM) reveals the highly crystalline nature of Ag2S
QDs with lattice fringes corresponding to d(1̅ 21) of 0.26 nm, with the QDs appearing to be
polycrystalline with some degree of crystal strain in most particles (Figure 5.1f). A Ag/S
ratio of ca. 2.06 was obtained from ICP-OES. This agrees well with the XPS
quantification which gives a Ag/S ratio of ca. 2.2, revealing the expected chemical
composition of QDs. Moreover, the monodisperse Ag2S QDs can be synthetically
reproduced with identical size and size distribution over multiple batches (Figure 5.1c).
The growth of well-defined QDs may be attributed to the steric hindrance provided by the
long alkyl chains at the N-substituents of the NHC ligands. In comparison, bulk Ag2S was
immediately obtained if (TMS)2S was reacted with AgNO3 without any organic ligands,
indicating that the NHC ligands are, not surprisingly, playing a critical role in QD
formation. Moreover, as compared to the conventional organic ligands, such as oleic acid
or oleylamine (Figure 5.3), the NHC ligands also provide superior competence in
preparing monodisperse Ag2S QDs under ambient conditions.
131

Figure 5.1 Characterization of the NHC-Ag 2S QDs after a 60 min reaction: (a) powder XRD
pattern; (b) UV-vis-NIR absorption spectrum; (c) size and size distribution (300 counts for each
sample) from three different synthetic batches; (d f) TEM micrographs. Inset in (f) is an HRTEM
image showing d
(1
̅
21)
= 0.26 nm.

Figure 5.2 (a,b): UV-vis-NIR absorption spectra of in-situ reaction of Ag 2S QDs (from 1-b-
AgBr) and Cu 2–xS (from 1-b-CuBr), respectively. (c e) TEM micrographs of Ag 2S QDs at
various time points. Size analysis reveals d = 7.8  0.8 nm (5 min), to 8.2  0.7 nm (30 min), and
10.3  0.6 nm (60 min) (300 counts for each). (f h) TEM micrographs of Cu 2–xS QDs at different
132
time points. Size analysis presents d = 6.8  1.0 nm (5 min), to 8.4  1.2 nm (30 min), and 8.8 
0.8 nm (60 min) (300 counts for each).

Figure 5.3 (a,b) TEM micrographs of Ag 2S QDs prepared from Ag(oleate), and
AgBr(oleylamine) under ambient conditions, respectively.

5.3.2. The Role of NHC Ligands: Tailoring the Steric and Electronic Properties
To further understand the role and examine the potential of NHC ligands in
controlling QD morphology, we subsequently varied the alkyl chain length of the N-
substituents, aiming to study the morphology dependence on the steric bulk of the NHC
ligands while keeping the electronic properties the same. Four additional benzimidazole-
based NHC-AgBr complexes with varying n-alkyl chain lengths at the N-substituents
(Figure 5.4), 2-b-AgBr (R = C10H21), 3-b-AgBr (R = C8H17), 4-b-AgBr (R = C6H13), and
5-b-AgBr (R = C2H5), were prepared and Ag2S QDs were synthesized via the same
reaction between j-b-AgBr (j = 1 5) and (TMS)2S at room temperature for 1 h. Powder
XRD results reveal that all the NHC-AgBr precursors give phase-pure, monoclinic Ag2S
133

Figure 5.4 Structures of benzimidazole- and imidazole-based NHC-AgBr precursors.

(Figure 5.5). However, compared to complex 1-b-AgBr, which only generates colloidal
Ag2S QDs, complexes 2-b-AgBr, 3-b-AgBr, and 4-b-AgBr all yield a mixture of
colloidal Ag2S QDs and aggregated precipitates of Ag2S; while complex 5-b-AgBr
results in only bulk Ag2S. The purified Ag2S QDs prepared from complexes 2-b-AgBr,
3-b-AgBr, and 4-b-AgBr possess a mean diameter of 9.7  0.6, 9.2  1.0, and 9.6  1.0
nm, respectively (Figure 5.6, Table 5.1). It is thus concluded that the size of Ag2S
nanocrystals does not depend strongly on the steric properties of the N-substituents on the
NHC ligands, with the exception of the shortest ethyl group (R = C 2H5) not being able to
arrest growth and giving bulk Ag2S. The weak dependence of Ag2S QD size on the alkyl
chain length may be explained by the fact that n-alkyl groups on the N-substituents can
fold away from the QD surface, and thus do not provide additional steric hindrance after
a certain chain length. The aggregates formed from j-b-AgBr (j = 2 4) likely result from
poorer colloidal stabilization with shorter chain lengths as compared to 1-b-AgBr. These
134
results suggest that long alkyl chains (C > 2) at the N-substituents of the NHC ligands
provide enough steric hindrance to arrest QD growth and provide varying degrees of
colloidal stabilization; however, the Ag2S QD size and size distribution do not correlate
with alkyl chain length on the N-substituents of the NHC ligands.

Figure 5.5 XRD patterns of the purified products prepared from different benzimidazole-based
NHC-AgBr complexes (j-b-AgBr, j = 1 5) (a), and imidazole-based NHC-AgBr complexes (j-i-
AgBr, j = 1 3) (b). All of the products are phase-pure monoclinic Ag 2S (PDF#00-014-0072).

Figure 5.6 (a d) TEM micrographs of Ag 2S QDs prepared from 1-b-AgBr (10.3  0.6 nm), 2-b-
AgBr (9.7  0.6 nm), 3-b-AgBr (9.2  1.0 nm), and 4-b-AgBr (9.6  1.0 nm) complex,
respectively. The insets on the upper right corner of each micrograph are photos of each solution
mixture after a 1 h reaction. No precipitates were observed in the reaction using 1-b-AgBr
complex, while black solids were observed from solutions using 2-b-AgBr, 3-b-AgBr, and 4-b-
AgBr complex. More aggregates were observed as the length of N-alkyl chains decrease (from 2-
b-AgBr, 3-b-AgBr, to 4-b-AgBr). The insets on the upper left corner of (c) and (d) were
representative TEM images of corresponding Ag2S QDs, showing aggregates upon synthesized.

135
Table 5.1 Summary of Ag 2S products from different NHC-AgBr precursors
precursor R TEP (cm
1
)
a
d (nm)
products
1-b-AgBr C 14H 29 2057 10.3  0.6 colloidal QDs
2-b-AgBr C 10H 21 2057 9.7  0.6 colloidal QDs with aggregates
3-b-AgBr C 8H 17 2057 9.2  1.0 colloidal QDs with aggregates
4-b-AgBr C 6H 13 2057 9.6  1.0 colloidal QDs with aggregates
5-b-AgBr C 2H 5 2057 N/A bulk precipitate
1-i-AgBr (4,5-H) C 14H 29 2054 7.9  0.5 colloidal QDs
2-i-AgBr (4,5-Ph) C 14H 29 2048 5.6  0.4 colloidal QDs
3-i-AgBr (4,5-Cl) C 14H 29 2059 10.5  1.6 colloidal QDs
a
TEP values are estimated based on calculated results on the corresponding NHC ligands with
methyl groups at both N-substituents.
43-44

Given the lack of positive correlation between the NHC ligand steric properties
with resulting QD size, we sought to investigate the effect of tuning the electron-donating
properties of the NHC ligands. Imidazole-based NHC-AgBr complexes, 1-i-AgBr (4,5-
H), 2-i-AgBr (4,5-Ph), and 3-i-AgBr (4,5-Cl) (Figure 5.4) with significantly different
electron-donating properties, were synthesized. In these complexes, the R groups at the
N-substituents are all C14H29, ensuring an identical steric parameter for these NHC
ligands. Following the same reaction protocol, phase-pure Ag2S was obtained from these
imidazole-based NHC-AgBr precursors, as identified by powder XRD. Interestingly, the
mean diameter of the resulting Ag2S QDs displays a linear correlation with the Tolman
electronic parameter (TEP)
45-46
of the ligand, which is a commonly employed merit of
electron-donating ability.
47
We estimated the TEP values of these synthesized NHC
136
ligands based on the calculated results for the corresponding NHC ligands with methyl
groups at both N-substituents.
43,44
This is a fair estimation, as the size of R groups has
minimal effects on the σ-donating ability
44,48
and all the j-i-AgBr (j = 1 3) NHC ligands
possess the same C14H29 groups at the N-substituents. The more electron-donating NHC
ligands (with lower TEP values) yield smaller Ag2S QDs with a more monodisperse size
distribution under otherwise identical reaction conditions. Both 1-i-AgBr (4,5-H) and 2-i-
AgBr (4,5-Ph) give Ag2S QDs with excellent size distribution ( /d = 6%), while 3-i-
AgBr (4,5-Cl) results in very poor control of size uniformity (Figure 5.7). This may be
attributed to the slower reaction kinetics between the initial nanocrystal nuclei with
monomers in 1-i-AgBr (4,5-H) and 2-i-AgBr (4,5-Ph).
35
Since the NHC ligand is an
almost pure σ-donor in NHC-coinage metal complexes, as revealed previously by Ghosh
and co-workers,
49
the more electron donating NHC ligands would result in stronger
NHC-Ag bonds. On the other hand, the more electron withdrawing NHC ligands lead to a
weaker NHC-Ag bond,
35
giving poorer control over QD size uniformity. Interestingly,
local QD superstructures are observed for the Ag2S QDs prepared from 2-i-AgBr (4,5-
Ph) (inset in Figure 5.7b), which additionally confirms the ability of this synthetic
approach to generate uniform, monodisperse Ag2S QDs.

137

Figure 5.7 (a c) TEM micrographs of Ag 2S QDs prepared from imidazole-based NHC-AgBr
precursors: 1-i-AgBr (4,5-H) , 2-i-AgBr (4,5-Ph), and 3-i-AgBr (4,5-Cl), respectively. The insets
on the upper left corner of each micrograph are the size and size distributions of the
corresponding QDs (300 counts each). The inset on the upper right corner of (b) is a TEM image
showing local superstructures of Ag 2S QDs prepared from 2-i-AgBr (4,5-Ph). (d) Plot of Ag 2S
QD mean diameter as a function of calculated TEP
43,44
of the NHC ligands.

5.3.3. Characterization of Surface NHC Ligands
The as-prepared Ag2S QDs from complex 1-b-AgBr exhibit excellent colloidal stability
in non-polar organic solvents (e.g., toluene, tetrachloroethylene), with no noticeable
precipitates after six months. As reported in our previous paper,
39
such colloidal stability
suggests a covalent interaction between the neutral NHC ligands and surface Ag( )
cations.
1
H NMR was first applied to confirm the presence of NHC ligands on the QD
surface. The same set of proton resonances was observed for the Ag2S QDs as for the
molecular NHC-AgBr precursor (Figure 5.8a), with significantly broadened peaks in the
case of the colloidal QDs. The broadening results from a slower tumbling rate of the
surface ligands, therefore indicating coordination of the NHC ligands to the Ag2S QD
surface.
50
The identity of the surface bound NHC ligands was also investigated by a
138
1
H 
13
C HSQC experiment, showing almost identical HSQC cross peaks for the Ag2S
QDs and the NHC-AgBr complex (Figure 5.8bc).

Figure 5.8 Characterization of surface NHC ligands on Ag 2S QDs. (a)
1
H NMR spectra of NHC-
AgBr and NHC-Ag 2S QDs in CD 2Cl 2 (from 1-b-AgBr). Resonances from 0.5 to 8 ppm are
assigned accordingly. Solvent impurities are indicated by  (toluene), * (CH 2Cl 2) and  (H 2O).
(b,c)
1
H 
13
C HSQC spectra of NHC-AgBr and a colloidal suspension of Ag 2S QDs in CD 2Cl 2,
respectively. (d f) High-resolution XPS spectra of Ag 3d, S 2p, and N 1s in Ag 2S QDs,
respectively.

NHC binding to the QD surface was further corroborated by XPS, which
displayed a N 1s peak at a binding energy of 401.1 eV (Figure 5.8f). Furthermore, XPS
revealed peaks at the expected binding energies for Ag 3d (3d3/2 at 373.5 eV and 3d5/2 at
367.5 eV) and S 2p (2p1/2 at 161.9 eV and 2p3/2 at 160.8 eV), consistent with the
composition and expected oxidation states in the Ag2S QDs.
51
It is also worth noting that
XPS shows that Br is present on the surface of the Ag2S QDs (Figure 5.9). The atomic
Br/Ag ratio was quantified by XPS to be ~13% for purified Ag2S QDs and therefore, the
139
surface is not entirely shelled with Br. Additionally, the presence of Br is not associated
with protonated benzimidazolium bromide as a proton resonance at ~11.6 ppm was not
observed in the
1
H NMR spectrum of the purified QDs. Thus, we attributed the observed
Br to unreacted NHC-AgBr precursors associated with the surface, or surface passivation
by Br
–
to charge balance excess Ag
+
cations.

Figure 5.9 High-resolution XPS spectra of NHC-AgBr (1-b-AgBr, a c), NHC-Ag2S
QDs (d g), and NHC-Cu2S QDs (h k). The absence of the strong Cu
2+
satellite peaks (at
942 eV and 962 eV) in (h) proves that the oxidation state of Cu2–xS NCs is mostly Cu
+
.

140
FT-IR spectra display characteristic bands for imidazolium ring stretching at 1560
cm
–1
(s, sym) and 1467 cm
–1
(s, sym)
52
in both the NHC-AgBr precursor and Ag2S QDs
(Figure 5.10a). The quantity of NHC ligands on the Ag2S QDs can be estimated from
TGA data. A single 7.3% mass loss event at 250 300 ˚C is assigned to the decomposition
and/or loss of NHC ligands on the surface of the Ag2S QDs (Figure 5.10b),
30,53
and the
number of NHC ligands is calculated to be 346 per nanocrystal on average (for d = 10
nm), giving a surface NHC density of ~1 NHC/nm
2
. Such surface density of NHC
ligands agrees well with our previous report.
39
More importantly, it was revealed
previously that the long N-alkyl chains point outwardly due to the steric hindrance and
the distance between the end corners of bended tetradecyl groups is about ~1 nm.
26

Therefore, a surface coverage of 1 NHC per nm
2
(1 nm  1 nm) seems highly feasible.

Figure 5.10 (a) FT-IR spectra and (b) TGA traces of NHC-AgBr (1-b-AgBr), NHC-Ag2S,
and NHC-Cu2–xS QDs.

5.3.4. Synthesis of Colloidal Cu2 –xS, Ag2Se, and Cu2 –xSe QDs
To demonstrate the generality of this synthetic approach, the same chemistry was
extended to copper sulfide (Cu2–xS) and silver and copper selenide (Ag2Se, Cu2–xSe) QDs
141
using the 1-b-MBr precursor (M = Ag, Cu). Following the same procedure, Cu2–xS
nanocrystals can also be obtained at room temperature. Similar reaction kinetics (i.e., fast
nucleation and slower nanocrystal growth) are observed by UV-vis-NIR absorption and
TEM analysis (Figure 5.2). Powder XRD reveals that the as-synthesized product
crystallizes as phase-pure, cubic digenite Cu1.8S (PDF no. 00-056-1256, Figure 5.11a). It
was found that the Cu2–xS phase does not change with different nominal sulfur
stoichiometry; for example, adding (TMS)2S with 1:1 (Cu:S) ratio yields the same
digenite Cu1.8S phase instead of CuS, suggesting that no oxidation/reduction reactions are
involved in the reaction. The UV-vis-NIR absorption spectrum displays a characteristic
localized surface plasmon resonance (LSPR) peak at ~1412 nm (Figure 5.11b). This is in
line with the nonstoichiometry nature of Cu2 xS nanocrystals, as revealed by XPS
quantification giving a Cu:S ratio of ca. 1.57. TEM micrographs show a mostly spherical
morphology in the resulting Cu1.8S nanocrystals, with a minority population possessing
anisotropic shapes such as cubes, rods, and triangles after a 60 min reaction (Figure
5.11c-e). The average size of the as-synthesized Cu1.8S nanocrystals is 8.8  0.8 nm ( /d
= 9%). An HRTEM image shown as the inset to Figure 5.11e reveals that Cu1.8S appears
to be single crystalline with a lattice spacing of 0.20 nm for the (220) planes.

142
Figure 5.11 Characterization of NHC-Cu 1.8S QDs after a 60 min reaction: (a) powder XRD
pattern; (b) UV-vis-NIR absorption spectrum exhibiting a broad LSPR peak; (c e) TEM
micrographs of Cu 1.8S QDs. Inset in (e) is an HRTEM image showing d
(220)
= 0.20 nm.

Colloidal Ag2Se and Cu1.8Se QDs can also be synthesized using the same NHC-
MBr precursors with (TBDMS)2Se.
41
(TBDMS)2Se is much easier to handle compared to
(TMS)2Se since it is a solid and less sensitive to both moisture and light. Both Ag2Se and
Cu1.8Se QDs can be obtained after a 60 min reaction between the NHC-MBr precursor
and (TBDMS)2Se. The UV-vis-NIR spectra exhibit featureless absorption into the NIR
for Ag2Se and a characteristic LSPR peak at ~1512 nm for the Cu1.8Se QDs (Figure
5.12b). Powder XRD reveals that the as-prepared Ag2Se QDs crystallize in a metastable
tetragonal phase (a = 0.706 nm and c = 0.498 nm)
54
instead of the thermodynamically
preferred orthorhombic phase (PDF no. 00-024-1041). It has been previously observed
that Ag2Se nanocrystals can be kinetically trapped in this metastable phase, which is not
observed in the bulk.
16,55
Our XRD pattern agrees very well with literature reports for the
tetragonal Ag2Se QDs.
10,16,55-57
TEM analysis reveals a spherical, monodisperse
morphology for the Ag2Se QDs with a mean diameter of 9.7  0.8 nm ( /d = 8%), and
HRTEM reveals a lattice spacing of 0.24 nm (Figure 5.12c), which is consistent with the
calculated d(220) = 0.25 nm of the tetragonal phase.
16
The as-prepared Cu2–xSe QDs
crystallize in the expected cubic Cu1.8Se phase (Figure 5.12a, PDF no. 01-071-6181) and
possess an quasispherical morphology with an average diameter of 10.1  1.2 nm ( /d =
12%) (Figure 5.12d). HRTEM analysis suggests that the QDs are single crystalline, with
observed lattice spacings of 0.33 nm corresponding to the (111) planes of Cu1.8Se. These
Ag2Se and Cu2–xSe QDs both possess excellent colloidal stability, with no observable
143
aggregates or decomposition (such as silver plating of the glass container, commonly
observed for Ag2Se QDs
16
) over six months. It is also worth noting that monodisperse
colloidal Ag2Se or Cu1.8Se QDs with sub-10 nm diameters have never been prepared
under ambient conditions to the best of our knowledge.

Figure 5.12 (a) Powder XRD patterns of Ag 2Se (peaks are assigned based on calculated
tetragonal Ag 2Se phase
16
) and Cu 1.8Se QDs; (b) UV-vis-NIR absorption spectra of Ag 2Se and
Cu 1.8Se QDs; (c,d) TEM micrographs of Ag 2Se and Cu 1.8Se QDs, respectively. Insets are the
HRTEM images of Ag 2Se and Cu 1.8Se QDs.

5.4. Experimental
5.4.1. Materials
Reagents and solvents were purchased from commercial sources and used as received,
unless otherwise stated. Benzimidazole (C7H6N2, 99%), imidazole (C3H4N2, 99%), 1-
bromotetradecane (C14H29Br), 1-bromodecane (C10H21Br), 1-bromooctane (C8H17Br), 1-
bromohexane (C6H13Br), 1-bromoethane (C2H5Br), silver( ) oxide (99+%), copper( )
oxide (99.9%), potassium carbonate (K2CO3, anhydrous, 99%), 1,4-dioxane,
bis(trimethylsilyl)sulfide ((TMS)2S, technical grade), tert-butyldimethylchlorosilane
144
(TBDMS-Cl, or
t
BuMe2Si-Cl, >98%), sodium (sticks, in mineral oil, 99%) and selenium
powder (-200 mesh, 99.999%) were purchased from Alfa Aesar. 1-octadecene (ODE,
technical grade, 90%), 4,5-diphenylimidazole (C15H12N2, 98%), 4,5-dichloroimidazole
(C3H2Cl2N2, >98%) were purchased from Sigma-Aldrich. ODE was dried under vacuum
overnight. Reactions involving air- or moisture-sensitive compounds were conducted
under a nitrogen atmosphere by using standard Schlenk techniques.

5.4.2. Synthesis of Benzimidazolium Salts, Imidazolium Salts, NHC-AgBr, and
NHC-CuBr Complexes
Benzimidazolium salts 1 5, benzimidazole-based NHC-AgBr (j-b-AgBr, j = 1 5) and
NHC-CuBr complexes, imidazole-based NHC-AgBr (j-i-AgBr, j = 1 3) complexes, and
(TBDMS)2Se were synthesized according to modified literature procedures.
26,41, 58, 59
The
following abbreviations are used: s = singlet, d = doublet, t = triplet, q = quadruplet, m =
multiplet, br = broad.

1,3-(Ditetradecyl)benzimidazolium bromide (1).
26
Benzimidazole (2.36 g, 20.0 mmol),
K2CO3 (2.76 g, 20.0 mmol), n-tetradecyl bromide (18 mL, 60
mmol) and CH3CN (20 mL) were added into a three-neck flask
and stirred at reflux (~85 ˚C) for 24 h. After the reaction, the solvent was removed under
reduced pressure, and the resulting solid was dissolved in CH2Cl2. The mixture was
filtered to remove the KBr precipitate. The filtrate was then concentrated under reduced
pressure. After, the residue was recrystallized from CH2Cl2/pentane and dried under
145
vacuum to yield a white solid (6.8 g, 57%).
1
H NMR (400 MHz, CDCl3)  11.56 (s, 1H),
7.71 7.63 (m, 4 H), 4.62 (t, J = 7.6 Hz, 4H), 2.10 2.00 (m, 4H), 1.45 1.2 (m, 44H), 0.87
(t, J = 7.1 Hz, 6H).
13
C NMR (125 MHz, CDCl3)  143.06, 131.46, 127.18, 113.18,
47.82, 32.04, 29.79, 29.76, 29.70, 29.68, 29.62, 29.51, 29.47, 29.17, 22.81, 14.24 
Scheme 5.2 General Synthetic Scheme

146
1,3-(Didecyl)benzimidazolium bromide (2). Benzimidazole (2.36 g, 20.0 mmol), K2CO3
(2.76 g, 20.0 mmol), n-decyl bromide (13 mL, 60 mmol) and
CH3CN (20 mL) were added into a three-neck flask and stirred at
reflux (~85 ˚C) for 24 h. After the reaction, the solvent was removed under reduced
pressure, followed by the dissolution in CH2Cl2. The mixture was filtered to remove the
KBr precipitate. The filtrate was then concentrated under reduced pressure. After, the
residue was recrystallized from CH2Cl2/pentane and dried under vacuum to yield a white
solid (3.4 g, 35%).
1
H NMR (500 MHz, CDCl3)  11.56 (s, 1H), 7.71 7.65 (m, 4 H),
4.62 (t, J = 7.49 Hz, 4H), 2.08 2.02 (m, 4H), 1.43 1.23 (m, 28H), 0.86 (t, J = 6.64 Hz,
6H).
13
C NMR (125 MHz, CDCl3)  142.94, 131.45, 127.19, 113.19, 47.82, 31.93, 29.66,
29.55, 29.48, 29.33, 29.15, 26.67, 22.75, 14.20.
1,3-(Dioctyl)benzimidazolium bromide (3). Benzimidazole (2.36 g, 20.0 mmol), K2CO3
(2.76 g, 20.0 mmol), n-octyl bromide (12 mL, 60 mmol) and
CH3CN (20 mL) were added into a three-neck flask and stirred
under reflux (~85 ˚C) for 24 h. After the reaction, the solvent was removed under
reduced pressure, followed by the dissolution in CH2Cl2. The mixture was filtered to
remove the KBr precipitate. The filtrate was then concentrated under reduced pressure.
After, the residue was recrystallized from CH2Cl2/pentane and dried under vacuum to
yield a white solid (2.5 g, 30%).
1
H NMR (500 MHz, CDCl3)  11.49 (s, 1H), 7.71 7.64
(m, 4 H), 4.62 (t, J = 7.52 Hz, 4H), 2.08 2.02 (m, 4H), 1.44 1.23 (m, 20H), 0.85 (t, J =
6.70 Hz, 6H).
13
C NMR (125 MHz, CDCl3)  142.85, 131.43, 127.20, 113.19, 47.81,
31.76, 29.63, 29.12, 29.08, 26.65, 22.65, 14.13.
147
1,3-(Dihexyl)benzimidazolium bromide (4). Benzimidazole (2.36 g, 20.0 mmol), K2CO3
(2.76 g, 20.0 mmol), n-hexyl bromide (5.6 mL, 40 mmol) and
CH3CN (20 mL) were added into a three-neck flask and stirred
under reflux (~85 ˚C) for 24 h. After the reaction, the solvent was removed under
reduced pressure, followed by the dissolution in CH2Cl2. The mixture was filtered to
remove the KBr precipitate. The filtrate was then concentrated under reduced pressure.
After, the residue was recrystallized from CH2Cl2/pentane and dried under vacuum to
yield a white solid (3.2 g, 44%).
1
H NMR (500 MHz, CDCl3)  11.60 (s, 1H), 7.71 7.65
(m, 4 H), 4.62 (t, J = 7.52 Hz, 4H), 2.08 2.02 (m, 4H), 1.42 1.33 (m, 12H), 0.87 (t, J =
7.52 Hz, 6H).
13
C NMR (125 MHz, CDCl3)  143.06, 131.46, 127.21, 113.18, 47.82,
31.27, 29.64, 26.35, 22.54, 14.06.
1,3-(Diethyl)benzimidazolium bromide (5). Benzimidazole (2.36 g, 20.0 mmol), K2CO3
(2.76 g, 20.0 mmol), n-ethyl bromide (3.0 mL, 40 mmol) and
CH3CN (20 mL) were added into a three-neck flask and stirred
under reflux (~85 ˚C) for 24 h. After the reaction, the solvent was removed under
reduced pressure, followed by the dissolution in CH2Cl2. The mixture was filtered to
remove the KBr precipitate. The filtrate was then concentrated under reduced pressure.
After, the residue was recrystallized from CH2Cl2/pentane and dried under vacuum to
yield a white solid (3.3 g, 66%).
1
H NMR (500 MHz, CDCl3)  11.28 (s, 1H), 7.75 7.62
(m, 4 H), 4.67 (q, J = 7.34 Hz, 4H), 1.70 (t, J = 7.36 Hz, 6H).
13
C NMR (125 MHz,
CDCl3)  142.04, 131.22, 127.25, 113.17, 42.99, 14.98.
148
NHC-AgBr (1-b-AgBr).
26
Ag2O (0.56 g, 2.4 mmol) was added to a solution of 1 (1.2 g,
2.0 mmol) in dried CH2Cl2 (40 mL). The mixture was refluxed
for 20 h, and excess Ag2O was filtered away. The filtrate was
concentrated under reduced pressure. After, the residue was
recrystallized from CH2Cl2/pentane and dried under vacuum to yield a light brown solid
(0.94 g, 74%).
1
H NMR (500 MHz, CD2Cl2)  7.51 (dd, J = 5.9 and 3.0 Hz, 2H), 7.43
(dd, J = 6.0 and 3.2 Hz, 2 H), 4.62 (t, J = 7.3 Hz, 4H), 2.10 1.9 (m, 4H), 1.4 1.1 (m,
44H), 0.87 (t, J = 6.7 Hz, 6H).
13
C NMR (125 MHz, CD2Cl2)  134.27, 124.41, 112.14,
50.14, 32.50, 30.81, 30.26, 30.23, 30.20, 30.12, 30.04, 29.93, 29.81, 27.40, 23.27, 14.45.
NHC-AgBr (2-b-AgBr). Ag2O (0.56 g, 2.4 mmol) was added to a solution of 2 (0.96 g,
2.0 mmol) in dried CH2Cl2 (40 mL). The mixture was refluxed
for 20 h, and excess Ag2O was filtered away. The filtrate was
concentrated under reduced pressure. After, the residue was
recrystallized from CH2Cl2/pentane and dried under vacuum to yield a white solid (0.65 g,
56%).
1
H NMR (500 MHz, CDCl3)  7.49 (dd, J = 6.2 and 3.1 Hz, 2H), 7.41 (dd, J = 6.1
and 3.1 Hz, 2 H), 4.39 (t, J = 7.3 Hz, 4H), 1.93 1.87 (m, 4H), 1.33 1.24 (m, 28H), 0.86
(t, J = 6.8 Hz, 6H).
13
C NMR (125 MHz, CDCl3)  133.79, 124.12, 111.66, 49.76, 31.97,
30.50, 26.61, 29.57, 29.39, 29.37, 27.00, 22.80, 14.24.
NHC-AgBr (3-b-AgBr). Ag2O (0.56 g, 2.4 mmol) was added to a solution of 3 (0.85 g,
2.0 mmol) in dried CH2Cl2 (40 mL). The mixture was refluxed
for 20 h, and excess Ag2O was filtered away. The filtrate was
149
concentrated under reduced pressure. After, the residue was recrystallized from
CH2Cl2/pentane and dried under vacuum to yield a light grey solid (0.62 g, 59%).
1
H
NMR (500 MHz, CDCl3)  7.48 (dd, J = 6.2 and 3.0 Hz, 2H), 7.41 (dd, J = 6.1 and 3.1
Hz, 2 H), 4.39 (t, J = 7.3 Hz, 4H), 1.93 1.87 (m, 4H), 1.36 1.24 (m, 20H), 0.86 (t, J =
6.5 Hz, 6H).
13
C NMR (125 MHz, CDCl3)  133.797, 124.13, 111.67, 49.77, 31.86,
30.50, 29.33, 29.23, 27.00, 22.73, 14.20.
NHC-AgBr (4-b-AgBr). Ag2O (0.56 g, 2.4 mmol) was added to a solution of 4 (0.73 g,
2.0 mmol) in dried CH2Cl2 (40 mL). The mixture was refluxed
for 20 h, and excess Ag2O was filtered away. The filtrate was
concentrated under reduced pressure. After, the residue was
recrystallized from CH2Cl2/pentane and dried under vacuum to yield a light grey solid
(0.47 g, 50%).
1
H NMR (500 MHz, CDCl3)  7.49 (dd, J = 6.5 and 3.4 Hz, 2H), 7.41 (dd,
J = 6.2 and 3.4 Hz, 2 H), 4.39 (t, J = 7.3 Hz, 4H), 1.94 1.86 (m, 4H), 1.33 1.31 (m, 12H),
0.87 (t, J = 7.3 Hz, 6H).
13
C NMR (125 MHz, CDCl3)  133.77, 124.14, 111.66, 49.76,
31.48, 30.45, 26.63, 22.60, 14.10.
NHC-AgBr (5-b-AgBr). Ag2O (0.56 g, 2.4 mmol) was added to a solution of 5 (0.51 g,
2.0 mmol) in dried CH2Cl2 (40 mL). The mixture was refluxed
for 20 h, and excess Ag2O was filtered away. The filtrate was
concentrated under reduced pressure. After, the residue was
recrystallized from CH2Cl2/pentane and dried under vacuum to yield a light grey solid
(0.33 g, 46%).
1
H NMR (500 MHz, CDCl3)  7.50 (dd, J = 6.1 and 3.1 Hz, 2H), 7.42 (dd,
150
J = 6.1 and 3.1 Hz, 2 H), 4.47 (q, J = 7.3 Hz, 4H), 1.53 (t, J = 7.3 Hz, 6H).
13
C NMR
(125 MHz, CD2Cl2)  133.48, 124.24, 111.58, 44.74, 16.12.
NHC-AgBr (1-i-AgBr).
58
Imidazole (2.0 g, 30 mmol), KOH (3.3 g, 60 mmol), n-
tetradecyl bromide (9.0 mL, 30 mmol) and CH3CN (17 mL) were
added into a three-neck flask under nitrogen and stirred under
reflux (~85 ˚C) for 20 h. After the reaction was finished, CH3CN was removed under
reduced pressure. The solid was then dissolved in CH2Cl2, washed with water (2  100
mL) and brine (100 mL), and dried by Na2SO4. An orange oil was obtained after the
liquid was concentrated by vacuum. Subsequently, n-tetradecyl bromide (9.0 mL, 30
mmol) and toluene (20 mL) were added to the product under nitrogen, and the mixture
was stirred under reflux for 48 h. The solvent was then evaporated to give a red oil (6).
Ag2O (0.56 g, 2.4 mmol) was added to a solution of 6 (1.2 g, 2.0 mmol) in dried CH2Cl2
(40 mL). The mixture was refluxed for 20 h before filtering off excess Ag2O. The CH2Cl2
was then removed under reduced pressure. An orange powder (1-i-AgBr, 0.26 g, 20%)
was obtained after adding excess acetone to the organic residue, and drying under
vacuum.
1
H NMR (500 MHz, CDCl3)  6.95 (s, 2H), 4.07 (t, J = 7.3 Hz, 4H), 1.82 1.77
(m, 4H), 1.29 1.25 (m, 44H), 0.88 (t, J = 6.8 Hz, 6H).
13
C NMR (125 MHz, CDCl3)
 120.78, 52.28, 32.08, 31.63, 29.84, 29.81, 29.77, 29.68, 29.59, 29.51, 29.30, 26.63,
22.85, 14.28.
NHC-AgBr (2-i-AgBr). 4,5-diphenylimidazole (3.3 g, 15 mmol), KOH (1.7 g, 30 mmol),
n-tetradecyl bromide (4.5 mL, 15 mmol) and CH3CN (9 mL)
were added into a three-neck flask under nitrogen and stirred
151
under reflux (~85 ˚C) for 20 h. After the reaction was finished, CH3CN was removed
under reduced pressure. The solid was then dissolved in CH2Cl2, washed with water (2 
100 mL) and brine (100 mL), and dried by Na2SO4. A white solid was obtained after the
liquid was concentrated by vacuum. Subsequently, n-tetradecyl bromide (4.5 mL, 15
mmol) and toluene (10 mL) were added to the product under nitrogen, and the mixture
was stirred under reflux for 48 h. The solvent was then evaporated to give a white solid
(7). Ag2O (0.56 g, 2.4 mmol) was added to a solution of 7 (1.4 g, 2.0 mmol) in dried
CH2Cl2 (40 mL). The mixture was refluxed for 20 h before filtering excess Ag2O. The
CH2Cl2 was then removed under reduced pressure. A white solid (2-i-AgBr, 0.78 g, 49%)
was obtained after adding excess acetone to the organic residue, and drying under
vacuum.
1
H NMR (500 MHz, CDCl3)  7.36 7.31 (m, 6H), 7.18 7.16 (m, 4H),
4.07 4.04 (t, 4H), 1.62 1.57 (m, 4H), 1.31 1.12 (m, 44H), 0.87 (t, J = 6.9 Hz, 6H).
13
C
NMR (125 MHz, CDCl3)  131.89, 130.55, 129.26, 128.89, 128.23, 50.01, 32.07, 31.77,
29.83, 29.80, 29.74, 29.62, 29.51, 29.44, 29.03, 26.53, 22.84, 14.27.
NHC-AgBr (3-i-AgBr). 4,5-dichloroimidazole (2.0 g, 15 mmol), KOH (1.7 g, 30 mmol),
n-tetradecyl bromide (4.5 mL, 15 mmol) and CH3CN (9 mL)
were added into a three-neck flask under nitrogen and stirred
under reflux (~85 ˚C) for 20 h. After the reaction was finished, CH3CN was removed
under reduced pressure. The solid was then dissolved in CH2Cl2, washed with water (2 
100 mL) and brine (100 mL), and dried by Na2SO4. An orange oil was obtained after the
liquid was concentrated by vacuum. Subsequently, n-tetradecyl bromide (4.5 mL, 15
mmol) and toluene (10 mL) were added to the product under nitrogen, and the mixture
152
was stirred under reflux for 48 h. The solvent was then evaporated to give a red oil (8).
Ag2O (0.56 g, 2.4 mmol) was added to a solution of 8 (1.2 g, 2.0 mmol) in dried CH2Cl2
(40 mL). The mixture was refluxed for 20 h before filtering off excess Ag2O. The CH2Cl2
was then removed under reduced pressure. A pale grey powder (3-i-AgBr, 0.28 g, 20%)
was obtained after adding excess acetone to the organic residue, and drying under
vacuum.
1
H NMR (600 MHz, CDCl3)  4.13 (t, 4H), 1.82 1.77 (m, 4H), 1.32 1.25 (m,
44H), 0.88 (t, J = 7.1 Hz, 6H).
13
C NMR (150 MHz, CDCl3)  117.17, 51.43, 32.07,
30.98, 29.83, 29.80, 29.76, 29.68, 29.56, 29.51, 29.27, 26.57, 22.84, 14.27.

NHC-CuBr (1-b-CuBr).
41
Cu2O (0.34 g, 2.4 mmol) was added to a solution of 1 (1.2 g,
2.0 mmol) in dried 1,4-dioxane (40 mL). The mixture was
refluxed for 20 h, and excess Cu2O was then filtered off. The
filtrate was concentrated under reduced pressure. The solid was
isolated by filtration, and washed with hexanes. After drying under vacuum, the title
compound was obtained as a brown solid (0.9 g, 74%).
1
H NMR (500 MHz, CD2Cl2)
 7.48 (dd, J = 6.0 and 3.0 Hz, 2H), 7.40 (dd, J = 6.0 and 3.0 Hz, 2 H), 4.62 (t, J = 7.2 Hz,
4H), 2.10 1.9 (m, 4H), 1.4 1.1 (m, 44H), 0.87 (t, J = 6.8 Hz, 6H).
13
C NMR (125 MHz,
CD2Cl2)  124.25, 111.94, 49.49, 32.50, 30.82, 30.26, 30.23, 30.20, 30.12, 30.04, 29.93,
29.78, 27.38, 23.27, 14.45.
(TBDMS)2Se.
59
0.62 g (27 mmol) Na, 1.07 g (13.0 mmol) Se powder, and 0.10 g (0.78
mmol) C10H8 were weighed into a three-neck Schlenk flask. 60 mL of dried THF was
added. The solution was refluxed under a nitrogen atmosphere for 4 h, followed by
cooling to room temperature. The dark solution was further cooled to 0 ˚C by an ice bath.
153
4.07 g (27.0 mmol) (
t
BuMe2Si)2Cl was added to the solution at 0 ˚C. The reaction
mixture was allowed to stir with warming to room temperature overnight, and residual
solids were filtered away. The filtrate was concentrated under reduced pressure and dried
under vacuum. A reddish solid product was then obtained (3 g, 75%).
1
H NMR (400
MHz, C6D6)  1.01 (s, 9H), 0.37 (s, 6H).
13
C NMR (125 MHz, C6D6)  26.47, 19.40, 0.38.

5.4.3. Control Experiments
Control experiment #1: Synthesis of Ag2S QDs using oleate ligands under ambient
conditions. Ag(oleate) was prepared by dissolving Ag2O (115 mg, 0.500 mmol) with
excess oleic acid (5 mL) at 100 ˚C for 2 h. A clear Ag(oleate) solution was obtained and
subsequently diluted with 15 mL of CH2Cl2 to give a 50 mM Ag(oleate) stock solution.
0.5 mL of (TMS)2S/ODE (0.1 M) was injected rapidly into 2 mL of the Ag(oleate) stock
solution. The reaction mixture was allowed to stir at room temperature for 1 h before
precipitation from excess acetone and redispersion in toluene.
Control experiment #2: Synthesis of Ag2S nanocrystals using oleylamine under ambient
conditions. AgBr(oleylamine) was prepared by dissolving AgBr (190 mg, 1.00 mmol)
with excess oleylamine (5 mL) at 120 ˚C for 3 h. A clear AgBr(oleylamine) solution was
obtained and subsequently diluted with 15 mL of toluene to give a 50 mM
AgBr(oleylamine) stock solution. 0.5 mL of (TMS)2S/ODE (0.1 M) was injected rapidly
into 2 mL of the AgBr(oleylamine) stock solution. The reaction mixture was allowed to
stir at room temperature for 1 h before precipitation from excess acetone and redispersion
in toluene.
154
5.4.4. Synthesis of Metal Sulfide (Ag2S, Cu2 xS) QDs
NHC-MBr (0.2 mmol) was first dissolved in 4 mL of dried CH2Cl2 under nitrogen
at room temperature. (TMS)2S was diluted with dried ODE to give a 0.1 M solution,
which was rapidly added with a 2:1 (M:S) stoichiometry into the stirred NHC-MBr
solution. ODE is used here as it is a well-known non-coordinating solvent. Caution:
(TMS)2S is extremely reactive and it can react with moisture in air rapidly producing H2S
gas! (TMS)2S/ODE solution was thus prepared using standard Schlenk techniques. The
reaction mixture changed from colorless to dark red within 1 5 min of addition
depending on various precursors, indicating the formation of metal sulfide QDs. Aliquots
were for UV-vis-NIR absorption and TEM analysis. After 60 min, the metal sulfide QDs
were purified by precipitation with excess acetone, and redispersed in toluene. Unreacted
precursors or agglomerated QDs were discarded by centrifugation.

5.4.5. Synthesis of Metal Selenide (Ag2Se, Cu2 xSe) QDs
NHC-MBr (0.2 mmol) was first dissolved in 4 mL of dried CH2Cl2 under nitrogen
at room temperature. (TBDMS)2Se was dissolved in dried THF to give a 0.1 M solution.
THF is used here since (TBDMS)2Se does not completely dissolve in ODE at room
temperature. Caution: (TBDMS)2Se is volatile and it can react with moisture in air
rapidly producing toxic H2Se gas! (TBDMS)2Se was stored at 20 C in a N2-flushed
glass vessel. (TBDMS)2Se /THF solution was prepared using standard Schlenk
techniques. The dissolved (TBDMS)2Se solution was added with a 2:1 (M:Se)
stoichiometry into the stirring NHC-MBr solution. The reaction mixture changed from
155
colorless to dark red within 1 5 min of addition depending on different precursors,
indicating the formation of metal selenide QDs. After 60 min, the metal selenide QDs
were purified by precipitation from excess acetone, and redispersed in toluene. Unreacted
precursors or agglomerated QDs were discarded by centrifugation.

5.4.6. Characterization
UV-vis-NIR spectroscopy was carried out on a Perkin-Elmer Lamba 950
spectrophotometer equipped with a 150 mm integrating sphere, using a quartz cuvette for
liquid samples. Thermogravimetric analysis (TGA) measurements were made on a TA
Instruments TGA Q50 instrument, using sample sizes of ca. 5 mg in an alumina crucible
under a flowing nitrogen atmosphere. TGA samples were prepared by drying the colloid
under flowing nitrogen at 80 ˚C for ~120 min, followed by lightly crushing the solid with
a spatula prior to analysis. FT-IR spectra were acquired from pressed KBr pellets on a
Bruker Vertex 80. Pressed pellets were made of dried metal chalcogenide QDs (~2 mg)
in a dry KBr matrix (~100 mg). Solution 1D and 2D
1
H,
13
C,
1
H 
13
C HSQC NMR
spectra were collected at ambient temperature on a Varian 500 spectrometer (500 MHz in
1
H) with chemical shifts represented in units of ppm. All spectra were referenced to the
residual solvent peaks (5.33 ppm for CD2Cl2, 7.26 ppm for CDCl3, 7.00 ppm for toluene-
d8). NMR samples were prepared by drying the QD colloid in a vacuum oven (~60 C)
and hexanes were added to disperse the QDs, separating them from NHC-metal
complexes since these precursors do not dissolve in hexanes. After drying the
QD/hexanes suspension in a vacuum oven, the QDs were redispersed in deuterated
156
solvents. QD dispersions were filtered through a 0.45 m filter before NMR
measurements. For molecular complexes, 32 scans with a delay time of 1 s were taken for
each sample, and the data are presented as averages of those scans. For QD samples, 64
scans with a delay time of 30 s were taken to allow complete relaxation between pulses
for
1
H spectra.
40
Typical NMR samples had a QD concentration in the range of 10 30
M. X-ray photoelectron spectroscopy (XPS) was conducted using a Kratos Axis Ultra
X-ray photoelectron spectrometer with an analyzer lens in hybrid mode. High-resolution
scans were performed using a monochromatic aluminum anode with an operating current
of 6 mA and voltage of 10 kV using a step size of 0.1 eV, a pass energy of 40 eV, and a
pressure range between 1 3  10
–8
torr. The binding energies for all spectra were
referenced to the C1s core level at 284.8 eV. Elemental analysis was conducted by a
Flash 2000 Elemental Analyzer (CHNS Elemental Analyzer). Samples (2 3 mg) were
prepared in tin crucibles. Powder X-ray diffraction (XRD) data was collected using a
Rigaku Ultima IV diffractometer in parallel beam geometry (2 mm beam width) using Cu
K  radiation (  = 1.54 Å). Samples were prepared by drop casting onto zero-diffraction,
single crystal Si substrates. Transmission electron microscopy (TEM) analysis was
performed on a JEOL JEM-2100 microscope at an operating voltage of 200 kV, equipped
with a Gatan Orius CCD camera. Samples for TEM analysis were prepared from dilute
purified QD suspensions deposited on 400 mesh carbon-coated copper grids (Ted Pella,
Inc.). Inductively coupled plasma-optical emission spectroscopy (ICP-OES) was
performed on the nanocrystal samples using a Thermo Scientific Icap 7000 series ICP-
OES. QDs were prepared by drying 10 mg of solid, digesting in concentrated HNO3, and
157
diluting with ca. 99 mL of 10% HNO3. The atomic ratios were calculated by averaging 3
independently synthesized QD samples.

5.5. Conclusions
In conclusion, we have presented a tunable room-temperature synthesis of coinage metal
chalcogenide QDs using NHC-MBr synthons. Phase-pure and highly monodisperse Ag2E
and Cu2–xE (E = S, Se) QDs with excellent reproducibility are obtained from the reaction
between NHC-MBr complexes and alkylsilyl chalcogenides. The mean size of the QDs
can be tailored by tuning the electronic parameter of the NHC ligands; however, the QD
size is not correlated with the linear n-alkyl chain length on the N-substituents of the
NHCs. A linear relationship is observed between the mean diameter of the Ag2S QDs and
the TEP value of the NHC ligand, demonstrating that more electron-donating NHC
ligands result in smaller QDs. This synthetic approach thus allows for QD size control
from the molecular precursor, rather than reaction time or temperature. With the rich
library of NHC structures and NHC-metal complexes, this synthetic route provides new
inspirations for: (1) novel surface functionalization of QDs with NHC ligands, and (2)
low temperature, tunable synthesis of other binary or ternary metal chalcogenide (e.g.,
SnS, SnSe, ZnS, ZnSe, CuInS2) QDs using NHC-metal complexes.

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163
Chapter 6. Accessing Metastable Wurtzite InP Nanocrystals Via N-
Heterocyclic Carbene Stabilized Plasmonic Cu3 xP*
*Manuscript in preparation.

6.1. Abstract
A colloidal approach to plasmonic Cu3 xP nanocrystals (NCs) is developed using N-
heterocyclc carbene (NHC) as the sole ligand. The NHC is found to be a useful ligand for
the stabilization of colloidal Cu3 xP NCs, which facilitate the preparation of metastable
wurtzite InP NCs through cation exchange. This demonstration of NHC-stabilized metal
phosphide NCs may provide further inspiration for accessing metastable nanomaterials
using this simple carbene chemistry.

6.2. Introduction
Colloidal metal phosphide nanocrystals (NCs) have emerged as an important class of
materials, as a result of their unique semiconducting, photovoltaic, catalytic, light-
emitting, plasmonic, and magnetic properties.
1 10
Self-doped colloidal copper phosphide
(Cu3 xP) NCs are particularly interesting as they possess both plasmonic and
semiconducting properties.
9,11 12
Different from traditional noble metals which display
limited tunability of their localized surface plasmon resonance (LSPR) in the visible
range, the LSPR band of Cu3 xP NCs has been tailored from near infrared (NIR) to mid-
infrared (MIR) region, offering promising alternatives especially for the pulsed laser
164
application in telecommunication.
12
Additionally, with the intrinsic p-doped character
and the observed bandgap of 1.3 1.4 eV, Cu3 xP NCs have been combined with n-type
CdS, showing appreciable photovoltaic efficiencies as a semiconducting absorber.
9
The
hexagonal Cu3 xP structure bears a robust anion framework, which serves as a facile
platform to access other metal phosphide NCs through cation exchange.
11

Despite the vast potential of Cu3 xP NCs, colloidal synthetic methods towards
well-defined Cu3 xP NCs remain underdeveloped. Early reports utilize trioctylphosphine
(TOP) to phosphorize copper NCs under relatively high reaction temperatures (i.e.,
320380 C).
11,13
Recently, more reactive phosphine molecules such as phosphine gas
(PH3) and tris(trimethylsilyl)phosphine ((TMS)3P) were introduced as the phosphide
source, which helps to conduct the reaction at moderate reaction temperatures (120 250
C).
9,12
Interestingly, it was found that the presence of those long-chain phosphine
molecules is very critical, as they act as a coordinating ligand (e.g., TOP), or as a
coordinating solvent (e.g., trioctylphosphine oxide, TOPO). When solely oleylamine is
used as the solvent under otherwise identical reaction conditions and no phosphine
ligands (i.e., TOP, TOPO) are involved in the reaction, only Cu NCs were produced.
12

Evidently, the neutral L-type phosphine ligands create a coordination sphere that can
manipulate the precursor reactivity and reaction kinetics, impart colloidal stability, and
tune the LSPR frequency of the resulting Cu3 xP NCs.
9,12
However, it remains
challenging to replace hazardous phosphine ligands in the synthesis of high-quality
colloidal Cu3 xP NCs. Hence, developing more advanced synthetic methods using
alternative ligands towards colloidal Cu3 xP NCs is of great importance, and requires a
165
better understanding of the coordination chemistry involved from molecules to NCs.
N-heterocyclic carbenes (NHCs) are an alternative neutral ligand set that forms
stronger coordination bonds with coinage metals than phosphine.
14
They have also been
applied to stabilize coinage-metal-based molecular complexes,
14
nanoclusters,
15
metal
NCs,
16
and metal chalcogenide NCs,
17 18
offering an excellent synthetic toolset to explore
the molecule to NC evolution. For instance, we have demonstrated that the size of Ag2S
NCs correlates linearly with the electron donating property (i.e., Tolman electronic
parameter (TEP)) of the NHC ligands, with higher electron donating NHCs producing
smaller Ag2S NCs.
17
However, the analogous NHC-stabilized metal phosphide NCs have
yet to be developed, and the synthetic potential of NHC ligands for metal phosphide NCs
has not been examined, although NHC-stabilized nanoscopic AgP clusters have been
isolated.
19
Herein, we report the first synthesis of colloidal plasmonic Cu3 xP NCs using
NHC as the sole ligand in a non-coordinating solvent, which can further serve as a parent
platform to access metastable wurtzite InP NCs.
6.3. Results and Discussion
Scheme 6.1 Synthesis of NHC-Cu 3xP NCs and the subsequent cation exchange reaction with
In
3+
/TOP.
166
Previously, we have shown that the metathesis reaction between NHC-metal
bromide (NHC-MBr, M = Ag, Cu) and alkylsilyl chalcogenide ((R3Si)2E, E = S, Se)
under ambient conditions yields monodisperse metal chalcogenide NCs, with the long-
chain NHC ligands arresting the NC growth.
17
We sought to extend the same chemistry
to prepare metal phosphide NCs by using (TMS)3P (Scheme 6.1). Initial attempts to
synthesize Cu3 xP NCs at room temperature resulted in an amorphous product, as
revealed by powder X-ray diffraction (XRD). Therefore, reaction temperatures were
raised to 250 C, and (TMS)3P was rapidly injected into a solution of NHC-CuBr in a
non-coordinating solvent (1-octadecene (ODE)). The as-synthesized Cu3 xP NCs were
confirmed by powder XRD to be phase-pure hexagonal Cu3 xP phase (PDF no. 00-002-
1263, Figure 6.1a). The vis-NIR absorption spectrum displays a characteristic LSPR band
at ~1360 nm (Figure 6.1b) as expected due to copper vacancies. Energy-dispersive X-ray
spectroscopy (EDX) analysis of the as-synthesized product gives a substoichiometric
formula of Cu2.76P, consistent with the copper-poor character (Figure 6.2a). Transmission
electron microscopy (TEM) images show a quasi-spherical morphology for the Cu3 xP
NCs with a mean diameter of 6.7  1.1 nm (300 counts), and high-resolution TEM
(HRTEM) reveals single crystalline Cu3 xP with a lattice spacing of 0.20 nm for the (300)
planes consistent with the hexagonal phase (Figure 6.1d). Such NHC-stabilized Cu3 xP
NCs are colloidally stable for days under inert atmosphere, which could be ascribed to
the surface-bound NHC ligands, as evidenced by the presence of N 1s peak in the XPS
data (Figure 6.4).
167
Figure 6.1 Spectroscopic characterizations of Cu 3xP and InP NCs after cation exchange: (a)
Powder XRD patterns; (b) vis-NIR absorption spectra; (c) Raman spectra. (d, e) TEM
micrographs of Cu 3 xP and InP NCs, respectively. Insets are the HRTEM images of Cu 3 xP and
InP NCs.

Figure 6.2 EDX spectrum of Cu 3xP (a) and InP (b) NCs, respectively. EDX data gives a Cu:P
ratio of 2.76 for Cu 3xP NCs, and a In:P ratio of 0.76 for InP NCs. EDX analysis also reveals
around 10% Br in InP NCs, which might be associated on the surface of InP NCs.

The as synthesized NHC-stabilized Cu3 xP NCs can then serve as an ideal
hexagonal platform to access metastable wurtzite InP NCs, a phase that is unusual to
168
obtain from traditional colloidal synthetic methods. Thus far, metastable wurzite InP
NCs have been observed mostly as a polytype in nanowires synthesized by molecular
beam epitaxy or chemical vapor deposition methods.
20 21
Manna’s group has recently
reported a colloidal approach to wurtize InP nanoplalets of 10 50 nm by cation exchange
with TOP-stabilized Cu3 xP nanoplalets.
9
Herein, we subsequently performed cation
exchange on NHC-stabilized Cu3 xP NCs through a reaction with InBr3/TOP solution
(In:Cu = 1:1) at 200 C. Powder XRD results indicate a complete cation exchange after
15 min, giving wurtzite InP NCs (ICSD card no. 180911, Figure 6.1a). The integrity of
the cation exchange reaction is confirmed by the absence of a LSPR peak and the
presence of InP absorption features with an onset at ~800 nm (Figure 6.1b). TEM
micrographs of InP NCs show the preserved morphology of the Cu 3 xP NCs, with a mean
diameter of 8.0  0.9 nm (300 counts, Figure 6.1d). It is noteworthy that the Cu3 xP NCs
were not isolated before reacting with In
3+
, thereby the slight increase of NC size might
be due to the surface reaction of In
3+
and (TMS)3P on the InP NCs. HRTEM displays
monocrystalline hexagonal InP, and atomic sketches represent [100], [010], and [1
̅
10]
lattice slabs of wurtzite InP (Figure 6.1d). Detailed analysis of these lattice fringes
reveals a slight compression along the ab plane (2%) as compared to the reference
wurtzite InP phase. The single domain of InP NCs without any visible defects suggests
the completion of the cation exchange reaction.
The efficacy of cation exchange reaction is further corroborated by Raman, EDX,
and X-ray photoelectron spectroscopy (XPS). Raman spectrum of InP NCs displays
characteristic Raman signals of the InP phase (Figure 6.3)
11,22
(transverse optical (TO)
169
phonon first order mode at 309 cm
1
; longitudinal optical (LO) phonon first order mode
at 347 cm
1
; and longitudinal optical (LO) phonon second order mode at 690 cm
1
) after
cation exchange, while the pristine Cu3 xP NCs do not exhibit any Raman features. EDX
elemental analysis of as-synthesized InP NCs gives an elemental composition of In0.76P
and only 2% of Cu (Figure 6.2b).
Figure 6.3 Raman spectra of Cu 3xP and InP NCs.

The high-resolution XPS spectrum of Cu 2p confirms the presence of Cu
+
in
Cu3 xP NCs with the expected binding energies of Cu 2p1/2 at 952.4 eV and 2p3/2 at 932.6
eV, and excludes any trace of Cu
2+
ions, in which case two Cu
2+
satellite peaks (at 942
eV and 962 eV) would be observed (Figure 6.4a). Additionally, the high-resolution XPS
spectrum of P 2p shows characteristic binding energies for P
3 
(2p1/2 at 129.2 eV, 2p3/2 at
128.4 eV), with a small amount of oxidized P species (e.g., POx) at 132.5 eV (Figure
6.4b). The poor air stability of Cu3 xP NCs is in agreement with literature reports
12
and
experimentally-observed results. However, when Cu3 xP NCs underwent cation exchange
and transformed to InP NCs, distinct In
3+
and P
3 
species (Figure 6.4 e,f) are identified
from high-resolution XPS spectra of In 3d (3d3/2 at 451.5 eV and 3d5/2 at 443.9 eV) and P
170
2p (2p1/2 at 129.3 eV, 2p3/2 at 128.5 eV). Moreover, oxidized P species are not observed
in the InP NCs. The as-synthesized InP NCs are colloidally stable for weeks. XPS
quantification from the survey spectrum gives a In to P ratio of 0.90. The residue of Cu
species can also be quantified to be ~2% with respect to In.
31
P NMR spectrum of a InP
NCs suspension shows two broad peaks at 44.4, and 26.8 ppm (d6-benzene, Figure 6.5a),
which are in close agreement with TOPO (42 ppm) and TOP ligands ( 33 ppm, d6-
benzene), respectively. Since the surface TOP ligands can be partially oxidized to TOPO
in air, the InP NCs are capped with both TOP and TOPO, and these reductive TOP
ligands are likely responsible for the improved air and colloidal stability of the InP NCs.

171
Figure 6.4 High-resolution XPS spectra of Cu 2p (a), P 2p (b) in Cu 3xP NCs and In 3d (c) and P
2p (d) in InP NCs. The absence of the strong Cu
2+
satellite peaks (at 942 eV and 962 eV) in (a)
indicates that the oxidation state of Cu in Cu 3–xP NCs is mostly Cu
+
. The additional peak at 133.7
eV in (b) proves the oxidation of P in Cu 3–xP NCs.

Figure 6.5
31
P NMR spectra of (a) a InP NCs suspension, (b) TOP, and (c) TOPO in d 6-benzene.
The broad peaks at 44.4, and 26.8 ppm can be assigned to TOPO (42 ppm) and TOP ligands
( 33 ppm), respectively. Additional sharp peaks from the InP suspension might be from
impurities in TOP/TOPO chemicals and they are not associated on the surface of InP NCs.

It has been recently revealed that such cation exchange reactions are mediated by
Cu vacancy diffusion,
11,23
with higher Cu vacancies accelerating cation exchange reaction
rates. However, our Cu3 xP intermediates were not isolated or exposed to air before
reacting with In
3+
, suggesting that Cu vacancies might be inherent during the synthesis.
In fact, it has been shown that the formation of copper vacancies is thermodynamically
favored.
11
Therefore, these hexagonal Cu3 xP NCs provide a facile and robust platform to
access other metal phosphide NCs through cation exchange reactions. Only zinc blende
InP NCs (Figure 6.6) with smaller diameters and lower reaction yield were produced
under otherwise identical reaction conditions when Cu3 xP intermediates were not
prepared first. Furthermore, we have shown that such cation exchange reactions are
associated with a ligand exchange process, providing additional chemical manipulation of
the NCs composition. The importance of surface coordinating ligands for cation
exchange reactions is often neglected; however, it also affects the thermodynamic
172
equilibrium of this exchanging process due to the significantly different binding affinities
of parent NCs and the foreign metals.
11
Thus, the utilization of these NHC-stabilized
Cu3 xP NCs might provide additional reaction parameters to tune the reactivity or cation-
exchange efficacy.

Figure 6.6 Characterization of zinc blende InP NCs: (a) powder XRD, (b) Uv-vis
absorption, (c) TEM micrograph. Size analysis of InP NCs gives a mean diameter of 3.13
 0.34 nm.

6.4. Experimental
6.4.1. General Considerations
Reagents and solvents were purchased from commercial sources and used as received,
unless otherwise stated. Benzimidazole (C7H6N2, 99%), 1-bromotetradecane (C14H29Br),
copper( ) oxide (99.9%), potassium carbonate (K2CO3, anhydrous, 99%),
tris(trimethylsilyl)phosphine, trioctylphosphine (TOP, 90%), indium tribromide and 1,4-
dioxane were purchased from Alfa Aesar. 1-octadecene (ODE, technical grade, 90%) was
purchased from Sigma-Aldrich. Reactions involving air- or moisture-sensitive
compounds were conducted under a nitrogen atmosphere by using standard Schlenk
techniques.
173
UV-vis-NIR spectroscopy was carried out on a Perkin-Elmer Lamba 950
spectrophotometer equipped with a 150 mm integrating sphere, using a quartz cuvette for
liquid samples.
Powder X-ray diffraction (XRD) data was collected using a Rigaku Ultima IV
diffractometer in parallel beam geometry (2 mm beam width) using Cu K  radiation (  =
1.54 Å). Samples were prepared by drop casting onto zero-diffraction, single crystal Si
substrates, followed by moderate heating to evaporate solvents.
X-ray photoelectron spectra (XPS) were obtained using a Kratos Axis Ultra X-ray
photoelectron spectrometer with an analyzer lens in hybrid mode. High resolution scans
were performed using a monochromatic aluminum anode with an operating current of 6
mA and voltage of 10 kV using a step size of 0.1 eV, a pass energy of 40 eV, and a
pressure range between 1 3  10
–8
torr. The binding energies for all spectra were
referenced to the C1s core level at 284.8 eV.
Transmission electron microscopy (TEM) analysis was performed on a JEOL
JEM-2100 microscope at an operating voltage of 200 kV, equipped with a Gatan Orius
CCD camera. Samples for TEM analysis were prepared from dilute purified nanocrystal
samples deposited on 400 mesh carbon-coated copper grids (Ted Pella, Inc.).
Solution
31
P NMR spectra were collected at ambient conditions on a Varian 500
spectrometer (500 MHz in
1
H) with chemical shifts represented in units of ppm. NMR
samples were prepared by drying nanocrystals (NCs) in a vacuum oven (~50 C),
followed by dispersing in deuterated solvents (d6-benzene). NC dispersions were then
filtered through a 0.45 m filter into a J Young tube for NMR measurements.

174
6.4.2. Synthesis and Characterization of Cu3 xP
In a typical Cu3 xP NCs synthesis, a solution of NHC-CuBr (195 mg, 0.3 mmol) in 6 mL
ODE was firstly prepared. The mixture was connected to vacuum and held for 1 h at 120
C to remove residue H2O. NHC-CuBr was observed to be completely dissolved in ODE
after 1 h stirring at 120 C. Subsequently, the temperature was raised to 250 C. (TMS)3P
was diluted with dried ODE (dried overnight under vacuum at 80 C) to give a 0.1 M
solution. 0.5 mL of (TMS)3P/ODE (with a 6:1 (Cu:P) stoichiometry) solution was rapidly
added into the NHC-CuBr solution. The reaction color changed from pale green to black
immediately upon addition. After 2 min, the reaction was quenched by removing the heat
and adding 10 mL of toluene. The Cu3 xP NCs were purified by precipitation with excess
acetone and redispersed in toluene or tetrachloroethylene (TCE). 

6.4.3. Synthesis of InP Via Cation Exchange Reaction
A solution of In
3+
/TOP was firstly prepared by dissolving InBr3 (106 mg, 0.3 mmol) in
TOP (2.2 mL), and dried ODE (5.3 mL) at 200 C for 1 h. The In
3+
/TOP solution was
then kept at room temperature under N2 atmosphere. Following the same reaction
procedures as Cu3 xP NCs synthesis, NHC-CuBr (195 mg, 0.3 mmol) was dissolved in
ODE and the reaction temperature was raised to 250 C before injecting (TMS)3P/ODE
solution. After reacting with (TMS)3P/ODE for 2 min, the In
3+
/TOP solution was added
into Cu3 xP NCs, and solution was allowed to react for 15 min at 200 C. The NCs were
then washed by dissolution in toluene followed by precipitation in excess acetone. The
175
InP NCs was finally dispersed in toluene or TCE. Reaction byproducts are not soluble in
toluene or TCE, and can be filtered to give a colloidal InP NCs.

6.4.4. Synthesis of Zinc Blende InP NCs
A solution of In
3+
/TOP was prepared by dissolving InBr3 (106 mg, 0.3 mmol) in TOP
(2.2 mL), and dried ODE (5.3 mL) at 200 C for 1 h. 0.5 mL of (TMS)3P/ODE (with a
6:1 (In:P) stoichiometry) solution was rapidly added into the In
3+
/TOP solution, and the
solution was stirred for 15 min at 200 C. The NCs were then precipitated with a mixture
of acetone/methanol (5:1), followed by dissolution in toluene or TCE.

6.5. Conclusions
We have reported a novel synthetic method towards colloidal Cu3 xP NCs using NHC as
the sole ligand in a non-coordinating solvent. The as-synthesized Cu3 xP NCs possess
characteristic NIR LSPR band, and crystalize as the expected hexagonal phase, which can
then serve as the parent platform to access colloidal InP NCs with the unusual metastable
wurtize phase via a cation exchange reaction. The integrity of both Cu 3 xP and InP NCs
are characterized thoroughly by various spectroscopic techniques. The cation exchange
reaction mediated by the NHC-stabilized Cu3 xP NCs is confirmed to be quantitative.
This successful demonstration of NHC-stabilized Cu3 xP NCs may inspire the synthetic
development of metal phosphide NCs or cation exchange reactions with carbene
chemistry.
176
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Abstract (if available)

Abstract Semiconductor nanocrystals, or quantum dots, are attractive functional materials for photovoltaics, photocatalysis, displays, and biomedical applications because of their uniquely tunable optoelectronic properties. The widespread implementation of these nanomaterials strongly relies on the sophisticated manipulation of their size, morphology, and surface functionality. Understanding nanocrystal surface chemistry and the nanocrystal-ligand interaction is particularly important as they could affect both the synthetic nanocrystal morphology and the optoelectronic properties. With this in mind, my Ph.D. work has been focused on designing new inorganic/organic surface ligands, and exploring the correlation between the surface chemistry and the charge transport properties of quantum-dots-based photovoltaic devices, as well as understanding how the nanocrystal-ligand interaction manipulates the resulting nanocrystal morphology. ❧ Given the low absorption efficiency as one of the limiting factors for current state-of-the-art hybrid polymer:nanocrystal solar cells, we sought to design the tandem and triple-junction devices for hybrid P3HT: CdSe nanocrystal solar cells. A combination of nanocrystalline ZnO and pH-neutral PEDOT: PSS was used as the interconnecting layer to serve as the Ohmic contact, and the thicknesses of subcells were optimized with the guidance of optical simulations. The tandem hybrid solar cells give a 2-fold increase in VOC over that of single-junction devices, and light absorption in the region of 480-600 nm is increased from 60% to 80% compared to the single-junction hybrid solar cells. Moreover, the average power conversion efficiency (PCE) was improved significantly from 2.0% (VOC = 0.57 V) in single-junction devices to 2.7% (champion 3.1%, VOC = 1.28 V) in tandem devices and 2.3% (VOC = 1.98 V) in triple-junction devices. The results clearly indicate that tandem/multijunction device geometries are an effective approach to harvest more light and improve PCE in hybrid BHJ solar cells. ❧ Another strategy to enhance absorption efficiency is to extend the semiconductor nanocrystal absorption threshold to near-infra (NIR) region which contains about 54% of solar energy. One of the most promising candidates is PbS quantum dots, in which the surface chemistry is much less understood as compared Cd-based quantum dots. For instance, it can be especially difficult to obtain colloidally stable suspensions of PbS nanocrystals ligand exchanged with small ligands, and many atomic ligands require dispersion in solvents that are incompatible with polymer solubility. We have developed a hybrid iodide-based surface chemistry, which affords air- and colloidally-stable PbS quantum dots in a polymer-compatible solvent. The modified PbS quantum dots offer a more efficient charge separation process when blended with the donor polymer (polymer poly[2,6-(4,4’-bis(2-ethylhexyl)dithieno[3,2-b:2’3’-d]silole)-alt-4,7-(2,1,3-benzothiadiazole)] (Si-PCPDTBT)), as evidenced by both photovoltaic performance and spectroscopic characterizations (i.e., transient absorption). The optimized Si-PCPDTBT: PbS hybrid solar cells give a broad spectral response from the visible region into the NIR, leading to a PCE of over 4% under AM 1.5G illumination. ❧ Additionally, we have expanded the boundary of semiconductor nanocrystal surface ligands from traditional carboxylic, phosphine, amine, thiol ligands to carbene ligands. We designed two synthetic approaches to N-heterocyclic carbene-capped semiconductor nanocrystals, which allows me to elucidate the energetic and dynamic nature of the obscure carbene-nanocrystal interface across different material types. It is revealed that these neutral L-type N-heterocyclic carbene ligands do provide strong, yet labile, binding to various coinage-metal-based semiconductor nanocrystals. As compared to coinage metal nanocrystals, these N-heterocyclic carbene ligands display stronger binding affinity to semiconductor nanocrystals without introducing any sub-bandgap trap states at the same time. This has further led to a more precise control of semiconductor nanocrystal morphology under ambient conditions via tailoring the electronic and steric properties of the carbene ligands. For instance, we have demonstrated that the size of Ag₂S NCs correlates linearly with the electron donating property of the NHC ligands (i.e., Tolman electronic parameter (TEP)), with higher electron donating NHCs producing smaller Ag₂S NCs. These interesting results have also intrigued us to expand the material platform to metal phosphide nanocrystals. We subsequently presented a facile synthetic route towards N-heterocyclic carbene stabilized colloidal plasmonic Cu₃₋ₓP nanocrystals, which can serve as a parent platform to access metastable wurtzite InP quantum dots that is unusual to obtain in other traditional colloidal synthetic methods.

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Asset Metadata

Creator Lu, Haipeng (author)

Core Title Expanding the library of surface ligands for semiconductor nanocrystal synthesis and photovoltaic applications

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 06/01/2017

Defense Date 05/08/2017

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag OAI-PMH Harvest,photovoltaics,semiconductor nanocrystal,surface chemistry,synthesis

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Brutchey, Richard (committee chair), Malmstadt, Noah (committee member), Thompson, Mark (committee member)

Creator Email hlv@usc.edu,hplv007@gmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c40-376060

Unique identifier UC11258352

Identifier etd-LuHaipeng-5341.pdf (filename),usctheses-c40-376060 (legacy record id)

Legacy Identifier etd-LuHaipeng-5341.pdf

Dmrecord 376060

Document Type Dissertation

Rights Lu, Haipeng

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

Repository Name University of Southern California Digital Library

Repository Location USC Digital Library, University of Southern California, University Park Campus MC 2810, 3434 South Grand Avenue, 2nd Floor, Los Angeles, California 90089-2810, USA