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Semiconductor inks for solution processed electronic thin‐films

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Semiconductor inks for solution processed electronic thin‐films

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Content SEMICONDUCTOR INKS FOR SOLUTION PROCESSED ELECTRONIC
THIN-FILMS

by

Jannise J. Buckley

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)

May 2016

Copyright 2016 Jannise J. Buckley

ii
Acknowledgements
I would not be writing this dissertation without the support from the Provost’s and
NSF fellowships or more importantly the encouragement and inspiration from my
advisors, coauthors and coworkers. Throughout my journey, they have given me unique
perspectives that have helped me learn and grow from both my successes and failures.
Without these special people and the lessons I have learned from them I would not be the
happy and confident woman I am today and for that, they will always have my sincerest
gratitude.
I would like to give special thanks to my advisor, Richard Brutchey. His
excitement for application driven inorganic chemistry and genuine concern for his
student’s scientific growth brought me to USC so I could specifically work under his
guidance. From the beginning he taught me to think independently, pushed me to secure
important scholarships, encouraged me to publish and present my research at conferences
around the world and most importantly, allowed me to pursue my own unconventional
career path. It has been a privilege to work in his lab alongside the members of the
Brutchey group who helped create a challenging, supportive and fun atmosphere to work
in.
Finally, I’d like to thank my family for their love and support. I thank my
grandparents, Joan and Bill Jordan, for their wisdom and encouragement and my parents,
Jill and Scott Buckley, for truly “being there”, in every sense of the phrase.

iii
Table of Contents

Acknowledgements ii

List of Tables vi

List of Figures vii

Abstract xiv

Chapter 1. Tin and Germanium Monochalcogenide IV –VI Semiconductor
Nanocrystals for Use in Solar Cells
1.1. Abstract
1.2. Introduction
1.3. General Structural Characteristics of Tin and Germanium
Monochalcogenides
1.4. Nanocrystal Synthesis and Characterization
1.4.1. Tin Sulfide Nanocrystals
1.4.2. Tin Selenide Nanocrystals
1.4.3. Tin Telluride Nanocrystals
1.4.4. Germanium Sulfide and Selenide Nanocrystals
1.5. Photovoltaic Device Applications
1.6. Conclusions
1.7. References

1
1
1

6
9
9
17
26
30
32
36
38

Chapter 2. Synthesis and Characterization of Ternary Sn
x
Ge
1 –x
Se
Nanocrystals
2.2. Introduction
2.3. Results and Discussion
2.3.1. Synthesis of Sn
x
Ge
1-x
Se Nanocrystals
2.3.2. Characterization of Ternary Sn
x
Ge
1-x
Se Nanocrystals
2.3.3. Formation Mechanism of Sn
x
Ge
1-x
Se Nanocrystals
2.3.4. Morphological Dependence on Composition
2.4. Experimental
2.4.1. Materials and Methods
2.5. Conclusions
2.6. References

49
49
51
51
53
62
64
65
65
66
67

iv
Chapter 3. Chalcogenol Ligand Toolbox for CdSe Nanocrystals and Their
Influence on Exciton Relaxation Pathways
3.1. Abstract
3.2. Introduction
3.3. Results and Discussion
3.3.1. Colloidal Ligand Exchange
3.3.2. NMR Characterization
3.3.3. Efficacy of Ligand Exchange
3.3.4. Photophysical Measurements
3.4. Experimental
3.4.1. Materials and Methods
3.5. Conclusions
3.6. References

71
71
72
75
75
78
85
91
101
101
105
106

Chapter 4. Ligand Exchange of Colloidal CdSe Nanocrystals with
Stibanates Derived from Sb
2
S
3
Dissolved in a Thiol-Amine
Mixture
4.1. Abstract
4.2. Introduction
4.3. Results and Discussion
4.3.1. Characterization of the Dissolved Sb
2
S
3
Species
4.3.2. Ligand Exchange Using Stibanates from Dissolved
Sb
2
S
3

4.3.3. Efficacy of Ligand Exchange
4.3.4. Nanocrystal Film Characterization
4.4. Experimental
4.4.1. Materials and Methods
4.5. Conclusions
4.6. References

112
112
112
115
115

121
125
131
135
135
139
140

Chapter 5. Dissolution of Sn, SnO and SnS in a Thiol –Amine Mixture:
Insights into the Identity of the Sn-Based Molecular Precursor
for Solution Processed SnS
5.1. Abstract
5.2. Introduction
5.3. Results and Discussion
5.3.1. Dissolution of Sn, SnO and SnS
5.3.2. Characterization of the Dissolved Sn Species
5.3.3. Control Reactions and Determination of Sn Oxidation

144
144
144
148
148
150

v
State
5.3.4. SnS Recovery Upon Annealing
5.4. Experimental
5.4.1. Materials and Methods
5.5. Conclusions
5.6. References

Bibliography

154
161
167
167
169
170

175

vi
List of Tables

Table 2.1.: Estimated compositions of Sn
x
Ge
1-x
Se determined using EDS 53

Table 2.2.: Rietveld Analysis of X-ray Diffraction Data of Ge
x
Sn
1 x
Se 58

Table 3.1.: Low temperature and high temperature mass loss percentages
from TGA
88

Table 3.2.: HOMO and LUMO levels of phenyl chalcogenols 97

Table 3.3.: Summary of fit constants for the PL decays 99

Table 4.1.: Raw data for elemental analysis on the stibanate ligands. 117

Table 4.2.: ICP-AES analysis of as-prepared and ligand exchanged CdSe
nanocrystals.

129

vii
List of Figures

Figure 1.1.: The GeS structure 7

Figure 1.2.: Calculated valence electron localization function for a
distorted GeS structure

8

Figure 1.3.: SEM images of SnS microcrystals

10

Figure 1.4.: TEM images of spherical SnS nanocrystals 13

Figure 1.5.: Photocurrent response of a thin film of SnS nanocrystals

14

Figure 1.6.: SEM image of 6-nm SnS nanocrystals

16

Figure 1.7.: Low- and high-resolution TEM image of SnSe nanocrystals

20

Figure 1.8.: TEM images and crystal structure of SnSe nanocrystals

23

Figure 1.9.: Tauc plot from UV-vis-NIR data for SnSe nanocrystals

24

Figure 1.10.: TEM images of 10-nm SnTe nanocrystals

27

Figure 1.11.: I-V curves for films of SnTe nanocrystals. 29

Figure 1.12.: TEM images of GeS and GeSe nanosheets

31

Figure 1.13.: Four-point I-V curve for dropcast films of GeSe nanosheets 34

viii

Figure 2.1.: XRD pattern of synthesized SnSe from control reaction
performed without the addition of HMDS
52

Figure 2.2.: XRD pattern of unknown GeSe phase

52

Figure 2.3.: Actual vs expected elemental composition of Sn
x
Ge
1-x
Se
nanocrystals

54

Figure 2.4.: Structural characterization of Sn
x
Ge
1 –x
Se nanocrystals 55

Figure 2.5.: Rietveld analysis of XRD patterns of Ge
x
Sn
1 x
Se samples 56

Figure 2.6.: Rietveld analysis of XRD patterns of Ge
x
Sn
1 x
Se samples

57

Figure 2.7.: Diffuse reflectance spectra for Sn
x
Ge
1 –x
Se nanocrystals

59

Figure 2.8.: Tauc plot of the linear regions from [F(R) hν]
1/2
as a function
of energy for the Sn
x
Ge
1-x
Se alloys

60

Figure 2.9.: Calculated indirect band gaps for Sn
x
Ge
1 –x
Se nanocrystals

60

Figure 2.10.: Timed aliquot XRD patterns taken durng the synthesis of
Sn
0.6
Ge
0.4
Se

62

Figure 2.11.: Lattice parameters of Sn
0.6
Ge
0.4
Se as a function of time 62

ix
Figure 2.12.: XRD patterns of the nanocrystal products taken when first
synthesizing SnSe and after injection of GeI
4
to the SnSe
nanocrystal product
63

Figure 2.13.: TEM images of Sn
x
Ge
1-x
Se nanocrystals 64

Figure 3.1.:
1
H NMR spectrum of the purified CdSe(NL) particles 79

Figure 3.2.:
1
H and
31
P NMR spectra of the digestion products from
CdSe(NL) nanocrystals

79

Figure 3.3.:
1
H and
77
Se NMR spectra of the reaction between Ph
2
Se
2
and
Ph
2
PH before and after the addition of CdSe(NL) nanocrystals

81

Figure 3.4.:
1
H NMR spectra monitoring the reaction between Ph
2
Se
2
and
Ph
2
PH
82

Figure 3.5.:
77
Se NMR spectra monitoring the reaction between Ph
2
Se
2
and
Ph
2
PH
83

Figure 3.6.:
31
P NMR spectra monitoring the reaction between Ph
2
Se
2
and
Ph
2
PH

84

Figure 3.7.: Magnified
1
H NMR spectrum of the reaction between Ph
2
Se
2
,
Ph
2
PH and CdSe(NL)

85

Figure 3.8.: Characterization of CdSe nanocrystals before and after ligand
exchange with phenylchalcogenols

87

Figure 3.9.: TGA traces of ligand exchange reactions with Bn
2
S
2
, Bn
2
Se
2
,
t
Bu
2
Se
2
, and
t
Bu
2
Te
2
precursors

87

x

Figure 3.10.: TEM images of CdSe(NL) and Ph
2
Te
2
exchanged particles

89

Figure 3.11.: Histograms of nanocrystal diameters before and after ligand
exchange with Ph
2
Te
2

90

Figure 3.12.: UV-vis of CdSe nanocrystals exchanged with various
ligands, tBuSeH, BnSeH and PhSeH

92

Figure 3.13.: Steady-state PL spectra of CdSe nanocrystals before and after
ligand exchange with phenylchalcogenols
95

Figure 3.14.: Control experiments performed to determine the effect of
phosphine on the PL intensity
96

Figure 3.15.: PL Plots showing the dependence of PL quenching on the
concentration of added Ph
2
Se
2
and Ph
2
PH

97

Figure 3.16.: Time-resolved PL spectra of CdSe nanocrystals before and
after ligand exchange with phenylchalcogenols

98
Figure 4.1.: Photograph before (left) and after (right) dissolution of 20 mg
bulk Sb
2
S
3
powder in 2 mL en and 0.05 mL ME

116

Figure 4.2.: A negative-ion ESI-MS spectrum of Sb
2
S
3
dissolved in 1:40
vol/vol mixture of en/ME at a concentration of 240 µg mL
–1

118
Figure 4.3.: Proposed chemical structures of the stibanate ligands derived
from Sb
2
S
3
dissolved in a solvent mixture of en/ME

120

Figure 4.4.: FT-IR spectrum of the dried stibanate ligands derived from
Sb
2
S
3
dissolved in a solvent mixture of en/ME
120

xi

Figure 4.5.: Raman spectrum of a solution of Sb
2
S
3
dissolved in en/ME

121

Figure 4.6.: Characterization of ligand exchange using zeta potential,
TGA and FT-IR

122
Figure 4.7.: Photograph showing the generality of the ligand exchange

124

Figure 4.8.:
13
C NMR spectra of the acid digestion products from the as-
prepared and ligand-exchanged CdSe nanocrystals

128
Figure 4.9.: High-resolution XPS spectrum of the Sb3d region of the
stibinate-capped CdSe nanocrystals after ligand exchange
and solution absorption and PL spectra of as-prepared and
ligand exchanged CdSe nanocrystals

130

Figure 4.10.: TEM micrographs of the as-prepared and ligand exchanged
CdSe nanocrystals

131
Figure 4.11.: Photocurrent response for ligand-exchanged CdSe
nanocrystal films

133
Figure 4.12.: Comparing photocurrent response for ligand exchanged
CdSe nanocrystal films heat treated to 300 °C and as-
prepared CdSe films heat-treated to 150 ºC

133
Figure 4.13.: FT-IR spectra of the dried, as-prepared and dried, ligand
exchanged CdSe nanocrystals

134

Figure 4.14.: XRD patterns comparing the ligand exchanged CdSe
nanocrystals before and after heating to 300 ºC

134

xii

Figure 5.1.: Spectroscopic characterization of Sn, SnO and SnS dissolved
in a 10:1 vol/vol 1,2-ethylenediamine / 1,2-ethanedithiol–
solvent mixture. A) Photograph, B)
119
Sn NMR spectra, C)
FT-IR spectra, and D) Raman spectra, of SnO (blue data),
SnS (purple data) and Sn (green data) dissolved in the en /
EDT solvent system

150
Figure 5.2.: Raman spectrum of the 10:1 vol/vol en/EDT solvent system

153

Figure 5.3.:
119
Sn NMR spectra of 1 in various solvents showing the
dependence of δ(
119
Sn) on Sn coordination number. The
green and purple shaded areas indicates the δ(
119
Sn) ranges
for four and six coordinate tin compounds, respectively

155
Figure 5.4.:
1
H (top) and
13
C (bottom) NMR spectra of 1 in CDCl
3

157
Figure 5.5.: Negative-ion ESI-MS spectrum of 1 dissolved in CHCl
3
in the
presence of en. The main ion cluster observed at m/z 362.9
is consistent with the formula [SnS
4
C
4
H
8
N
2
C
2
H
7
]
-

158
Figure 5.6.: Raman spectrum of 1 dissolved in en/EDT. The band at 178
cm
-1
is assigned to both the δ(N-Sn-S) and ν(Sn-N)

158

Figure 5.7.: Sn(II) vs Sn(IV) coordination geometry

159
Figure 5.8.:
119
Sn NMR spectra of SnCl2 2H
2
O and SnCl
4
5H
2
O
dissolved in the en/EDT solvent system

161

xiii
Figure 5.9.: Thermogravimetric, spectroscopic and optical
characterization of the recovered SnS. A) TGA/DSC traces
comparing the decomposition of 1 to the dried Sn, SnO and
SnS / en / EDT systems. B) FT-IR spectra of the SnO-ink
dried to 107 °C for 24 hours under vac and after annealing to
350 °C. C) Powder XRD patterns of the recovered SnS
obtained after annealing the dried Sn-inks to 350 °C under
flowing N
2
(g). D) Tauc plots showing the indirect band gaps
of the recovered SnS.

162

Figure 5.10.: Mass spectra of the evolved gases obtained from TGA/DSC
at ~315 ºC for all samples.

163
Figure 5.11.: Diffuse reflectance spectra of SnS obtained from annealing
the Sn, SnO and SnS inks at 350 ºC under N
2
(g).

165
Figure 5.12.: XRD patterns of the recovered 1 and SnCl
4
5H
2
O controls
from the en / EDT solvent system.

166

xiv
Abstract
Solution processing of inorganic semiconductors for the fabrication of affordable,
scalable electronic thin-films is emerging as a commercially viable alternative to
traditional high vacuum deposition methods. Using this approach, semiconductor thin
films can be printed onto a various substrates (e.g. plastics, glasses, or textiles) at low
temperatures using existing high-throughput technologies. In this work, we have
developed a number of strategies to synthesize, understand and enhance nanocrystalline
and molecular inks for the deposition of high-quality inorganic semiconducting films.
Compositionally controlled Sn
x
Ge
1 –x
Se nanocrystal inks were synthesized for the
first time as an environmentally sound alternative to lead- and cadmium-based
chalcogenides. The band gaps, lattice parameters, and morphologies of the alloys were
easily tuned via nanocrystal composition. Interestingly by alloying we showed facile
band gap tuning, which is otherwise difficult to achieve in layer structure IV –VI
semiconductors.
As synthesized, semiconductor nanocrystal inks are not ready for use in
nanocrystal-based devices due the presence of electrically insulating long-chain organic
surface ligands. While these native ligands ensure chemical and colloidal stability, they
must be removed in order to enhance the conductivity of the nanocrystalline inks for
electrically conductive films. With this in mind, we have developed a number of ligand
exchange techniques whereby the electronically insulating ligands on cadmium selenide
(CdSe) nanocrystals are replaced with smaller molecules to enhance the electron mobility
of the resulting thin-films. In one approach we installed small chalcogenol ligands via
the in situ reduction of R
2
E
2
(R =
t
Bu, Bn, Ph; E = S, Se, Te) by diphenylphosphine
(Ph
2
PH). The nanocrystal photophysics were investigated using time-resolved and low
temperature photoluminescence measurements. It was found that the phenylchalcogenol
ligands rapidly quench the photoluminescence by hole localization onto the ligand. In
another approach, molecular stibanates, derived from the dissolution of bulk Sb
2
S
3
in a
binary ethylenediamine and mercaptoethanol solvent mixture were employed as capping
ligands. Photoelectrochemical measurements on stibinate-capped CdSe nanocrystal films
showed that this novel ligand leads to a > 25-fold increase in photocurrent response
relative to as-prepared CdSe nanocrystal films.
Molecular inks are an attractive alternative to nanocrystalline based inks because
they do not require laborious synthesis, purification or surface tuning steps. In this work,
tin sulfide based inks were made by dissolving bulk tin precursors in an ethylenediamine
(en) / ethanedithiol (EDT) solvent mixture. Despite numerous publications utilizing this
solvent system, there is no direct information regarding the chemical identity of the
dissolved species. Bisethanedithiolate tin(II) was identified as the single tin species
present in all three (Sn, SnS and SnO) / en / EDT solutions, despite the various bulk

xv
precursor compositions (Sn, SnO and SnS) and oxidation states (Sn
0
and Sn
2+
). All three
inks can be converted to SnS using a mild annealing step (~350 ºC). Therefore, using a
simple “dissolve and recover” process, cheap bulk materials can be used as elemental
sources for high quality semiconducting inks and thin-films.

!
!
1
Chapter 1. Tin and Germanium Monochalcogenide IV–VI Semiconductor
Nanocrystals for Use in Solar Cells*
*Published in Nanoscale 2011, 3, 2399–2411.

1.1. Abstract
The incorporation of colloidal semiconductor nanocrystals into the photoabsorbant
material of photovoltaic devices may reduce the production costs of solar cells since
nanocrystals can be readily synthesized on a large scale and are solution processable.
While the lead chalcogenide IV–VI nanocrystals have been widely studied in a variety of
photovoltaic devices, concerns over the toxicity of lead have motivated the exploration of
less toxic materials. This has led to the exploration of tin and germanium
monochalcogenide IV–VI semiconductors, both of which are made up of earth abundant
elements and possess properties similar to the lead chalcogenides. This chapter highlights
recent efforts made towards achieving synthetic control over nanocrystal size and
morphology of the non-lead containing IV–VI monochalcogenides (i.e., SnS, SnSe,
SnTe, GeS and GeSe) and their application toward photovoltaic devices.

1.2. Introduction
The need for an affordable, secure, and sustainable energy landscape has motivated
governments and researchers across disciplines to explore alternative forms of energy,
such as solar power. The sun has the potential to supply the world energy demand of ca.
!
!
2

363 terawatt hours per day in just seconds, but the high cost of current silicon-based
photovoltaics (PVs) has been limiting.
1–3
Consequently, much work has focused on
decreasing processing costs through the development of organic- and polymer-based
PVs,
4,5
but these devices typically have the disadvantage of a narrow absorption range in
addition to poor thermal and environmental stability.
6,7
Other work has focused on
colloidal inorganic nanocrystals which share the synthetic advantages of scalability and
solution processibility with organic and polymeric materials, but with a potentially wider
absorption of the solar spectrum and superior transport properties.
8

The potential of utilizing colloidal semiconductor nanocrystals in PV devices has led to
the development of low-cost processing methods, such as spin-coating,
9,10
ink-jet
printing,
11,12
spray coating,
13
and drop casting,
14
at low temperatures and without the need
for expensive high vacuum processing methods. Solution-based synthetic methods can
now achieve precise control over the nanocrystal size and shape through the manipulation
of synthetic variables, such as reaction temperature, reaction time, reaction solvent, and
the ratio of capping ligand to precursors.
15–25
As a result, the band gap (E
g
) of the
resulting semiconductor nanocrystal can be tuned through control of the nanocrystal's
size and shape below the dimension of its Bohr exciton radius, which ultimately allows
for a fine level of control over the energy needed to inject an electron into the conduction
band of the material.
26
The lead chalcogenide IV–VI semiconductor nanocrystals (e.g.,
PbS and PbSe) have been explored as potential earth abundant active layers in PV
!
!
3
devices because of the high level of synthetic control that has been achieved in these
systems.
27–29

Lead chalcogenide nanocrystals were the first of the IV–VI class of semiconductor
nanocrystals to receive experimental interest as practical PV materials. Beginning in
1985–1986, Nozik and coworkers published the first reports on quantum confinement
effects in lead chalcogenides after observing size dependent shifts in the optical
absorption spectra of PbS and PbSe nanocrystals.
30,31
The strong quantum confinement
effects observed in lead chalcogenides originate from their relatively large Bohr exciton
radii, which range from 20 nm for PbS to 46 nm for PbSe.
32
Comparatively, other II–VI
and III–V classes of semiconductors (e.g., CdTe and GaAs) have smaller Bohr exciton
radii that are <15 nm.
32
An important consequence of the large Bohr exciton radii is the
ability to tailor band gap energies of the lead chalcogenides; for example, the band gap of
PbSe can be tuned from 1.2–0.5 eV by increasing the nanocrystal diameter from 2–15
nm, respectively.
33
The ease of size tunability on the nanoscale becomes specifically
useful in device design. Consequently, lead chalcogenide nanocrystals with controlled
size and morphology have been incorporated into to a variety of PV devices.
In 2008, Nozik et al. prepared an all-inorganic lead chalcogenide/aluminium Schottky
junction solar cell based on colloidally synthesized nanocrystal thin films.
34
It was found
that the most efficient devices resulted from thin films of PbSe nanocrystals that were 4
nm in diameter (E
g
= 0.9 eV). The PbSe-based device was fabricated from ITO/PbSe/Al
layers which gave large short-circuit photocurrents (J
SC
= 24.5 mA cm
−2
), low open-
!
!
4
circuit voltages (V
OC
= 239 mV), modest fill factors (FF = 0.41) and a spectrally
corrected AM1.5G power conversion efficiency of η
P
= 2.1%. The devices did not
require sintering or superlattice ordering for charge carrier transport; however, the device
was highly air sensitive and light absorption was hindered by thin active layers required
by the placement of the Schottky junction at the back contact. Although device
architectures based on Schottky junctions show promise for solar energy conversion, they
are typically limited by low open circuit voltages. In order to mitigate these limitations,
alternative device architectures such as p–n or p–i–n junctions may be necessary.
Moreover, third generation solar cells that might take advantage of multiple exciton
generation (MEG) in lead chalcogenide nanocrystals may be another way to improve
device efficiencies past current values through the generation of two or more excitons
from a single incident photon.
35–37

Despite the extremely promising results obtained thus far for lead chalcogenide-based
PVs, the toxicity of lead and the potential of lead exposure are considered a significant
threat to public health.
38,39
Lead is a poisonous substance known to affect most systems
in the body, including the production of red blood cells, the kidneys, and the central
nervous system.
40,41
Recent studies have even shown adverse effects (such as intellectual
deficits) in children with low bloodstream lead levels (BLLs < 10 µg dL
−1
).
42
In response
to lead's adverse physiological effects, the European Food Safety Authority (EFSA) has
recently decreased the tolerable exposure level to 0.50 µg kg
−1
bodyweight per day.
43

The decreasing threshold for BLLs observed over the past 20–30 years underscores the
!
!
5
importance of exploring alternative semiconductor materials that are less toxic, but have
otherwise similar properties to the lead chalcogenides. Tin(II) and germanium(II)
monochalcogenides are less toxic than the lead chalcogenides; and unlike lead, inorganic
tin compounds are not readily absorbed into the bloodstream through ingestion or
inhalation.
44
Although studies have cautioned over excessive exposure to germanium,
45
the toxicity of most germanium compounds is considered low.
46
The tin and germanium
monochalcogenides are made up of earth abundant elements,
47
and posses many other
properties similar to the lead chalcogenides, which make them particularly attractive
targets as the photoabsorbant material in nanocrystal-based PVs.
While the tin and germanium monochalcogenides (i.e., SnS, SnSe, SnTe, GeS, and GeSe)
appear to be promising replacements for the lead chalcogenides in PV applications, the
number of reports of tin and germanium monochalcogenide nanocrystals in the literature
over the past decade remains relatively small. Moreover, the same degree of synthetic
control over size and morphology that has been achieved in the lead chalcogenides has
not yet been realized in these related systems. This chapter highlights the recent work by
our group and others on the synthesis of high quality tin and germanium
monochalcogenide nanocrystals and their application towards PV devices. It will be seen
that while the degree of synthetic control over these systems is approaching that of the
lead chalcogenides, their application toward nanocrystal-based PV devices still remains
limited, albeit with a great deal of potential promise.

!
!
6
1.3. General Structural Characteristics of Tin and Germanium Monochalcogenides
The germanium and tin chalcogenides possess rather unique optical and electronic
properties resulting from the diverse compositions and structures of these materials. An
example of this diversity can be seen in the family of tin monochalcogenides where both
layered 2-D (e.g., SnS and SnSe) and 3-D (e.g., SnTe) crystal structures are observed.
This rich structural chemistry is partially a result of the Sn
2+
oxidation state in these
materials (vide infra), and can be further complicated by the ability of tin to access the
Sn
4+
oxidation state and possess coordination numbers in the solid state ranging between
2 and 9.
48,49

The two most common crystal structures observed in germanium and tin
monochalcogenide IV–VI semiconductors are the cubic NaCl (rock salt) and the
orthorhombic GeS structures. In bulk, the sulfides and selenides (i.e., GeS, GeSe, SnS,
and SnSe) possess an orthorhombic Pnma structure at low temperatures.
49,50
This
structure is best described as a highly distorted rock salt structure that is comprised of
zig-zag double layer planes of the metal monochalcogenide separated by a van der Waals
gap of ca. 1 Å (Fig. 1). The local arrangement about the Sn
2+
or Ge
2+
cations is that of a
distorted octahedron (coordination number = 6). In this structure, the cation–anion bond
angles deviate slightly from 90°, and there are three short and three long bonds about the
cation. This distortion can be primarily attributed to the Sn(5s) and Ge(4s) lone pairs for
Sn
2+
and Ge
2+
, respectively (Fig. 2).
51,52
The strong anisotropic properties observed in
SnS and SnSe, for example, have been correlated to this double layer structure, and it has
!
!
7
been observed that the electrical conductivity and Hall mobility at room temperature are
five to six times higher when measured along the layers instead of along the
crystallographic c axis.
53

Figure 1.1. The GeS structure. The yellow shaded atoms represent sulfur, while the aquamarine
shaded atoms represent germanium.

Tin sulfide (SnS) occurs naturally as herzenbergite in rare mineral form. At room
temperature, the thermodynamically preferred phase is that of α-SnS (a = 11.14 Å, b =
3.97 Å, and c = 4.34 Å). Above 600 °C, the Pnma structure converts into the
orthorhombic Cmcm β-SnS structure (a = 4.12 Å, b = 11.48 Å, and c = 4.17 Å).
54
Instead
of a metastable cubic phase,
55
it was found that SnS exhibits a pseudotetragonal crystal
structure (a = 11.55 Å, b = 4.12 Å, and c = 4.12 Å) with decreasing particle size.
56

Similarly, the thermodynamically preferred phase of α-SnSe is the orthorhombic Pnma
structure (a = 11.50 Å, b = 4.15 Å, and c = 4.45 Å), which undergoes a phase transition at
ca. 523 °C to the orthorhombic Cmcmβ-SnSe structure (a = 4.31 Å, b = 11.70 Å, and c =
4.31 Å).
54,57
At room temperature, β-SnTe possesses a cubic Fm3m structure (a = 6.32
Å), similar to the lead monochalcogenides, that is stable up to 727 °C at which point a
!
!
8
congruent melt occurs.
58
Solid-state studies have shown a low temperature cubic β-SnTe
to rhombohedral α-SnTe phase transition occurs at ca. −200 °C.
54

The thermodynamically preferred phase of GeS at room temperature is that of the
orthorhombic Pnma α-GeS structure (a = 10.47 Å, b = 3.64 Å, and c = 4.30 Å). Above
600 °C, the orthorhombic Pnma structure undergoes a phase transition to the hexagonal
β-GeS structure (a = 8.70 Å, c = 8.73 Å).
59
Similarly, the thermodynamically preferred
phase of GeSe at room temperature is also an orthorhombic Pnmaα-GeSe structure (a =
4.40 Å, b = 3.85 Å, and c = 10.82 Å).
60
This phase undergoes a phase transition at ca.
586 °C to the cubic Fm3m β-GeSe structure (a = 5.73 Å).
61

Figure 1.2. Calculated valence electron localization function for a distorted GeS structure.
Lighter colors signify regions of strong electron localization, while darker colors signify poorly
localized regions. The Ge(4s) lone pair is observed as an area of high electron localization at the
top right of the center cation. Reprinted with permission from ref. 64 (Copyright 2003 American
Physical Society).

The tin monochalcogenides possess an electron configuration of 4d
10
5s
2
5p
0
for Sn
2+
;
where the two Sn(5p) electrons are involved in bond formation, and the two Sn(5s)
!
!
9
electrons act as a lone pair.
49
The electron configuration for the corresponding
chalcogenide anion is ns
2
np
6
. The band structure for SnS and SnSe is such that the main
contribution to the top of the valence band is from the p orbitals of S
2−
or Se
2−
(with some
degree of hybridization with the cation s band), while the main contribution to the bottom
of the conduction band is from the empty p orbitals of Sn
2+
.
49,62–65
The Sn(5s) lone pair
does not participate in bonding to a great extent, but leads to a distortion from octahedral
geometry about the Sn
2+
cations (vide supra). Tight binding calculations predict that the
direct band gap decreases down the group for bulk tin monochalcogenides (E
g
= 2.1 eV
for SnS, 1.3 eV for SnSe and 1.1 eV for SnTe). Moreover, it has been calculated that
both bulk SnS and SnSe possess indirect band gaps that are close in energy to the direct
band gaps (E
g
= 1.5 eV for SnS and 0.9 eV for SnSe).
49
Similar to the tin
monochalcogenides, GeS and GeSe also possess closely placed direct and indirect band
gaps (Eg = 1.6–1.7 eV for GeS and Eg = 1.1–1.2 eV for GeSe) that overlap well with the
solar spectrum.
66,67

1.4. Nanocrystal Synthesis and Characterization
1.4.1. Tin Sulfide Nanocrystals
The bulk properties of SnS have been studied extensively and it has been found that the
material is stable under ambient conditions. Consistent with its layered crystal structure
(vide supra), SnS is an anisotropic native p-type semiconductor with reported hole
mobilities on the order of 90 cm
2
V
−1
s
−1
perpendicular to its c axis at 27 °C.
68,69
As a
!
!
10
result of its bulk transport properties and high absorption coefficient of ca. 10
4
cm
−1
, SnS
has promise as an absorber in PV devices.
70,71

Figure 1.3. (a) SEM images of SnS microcrystals synthesized from SnCl
2
and elemental sulfur.
(b) XRD pattern of the SnS product. (c) Cubic unit cell of zinc blende SnS viewed off the (100)
axis (top) and down the (111) axis (bottom). Reprinted with permission from ref. 77 (Copyright
2006 Wiley-VCH).
!
In an effort to explore the properties of SnS on the nanoscale, several solvothermal
methods have produced 0-D spherical particles, 1-D whiskers, and even 3-D tetrahedral
nano- and microcrystals.
72–77
In 2006, Greyson et al. synthesized SnS nano- and
microcrystals with a metastable zinc blende crystal structure which differs from the
!
!
11
thermodynamically preferred orthorhombic phase.
77,56
The SnS nano- and microcrystals
were prepared by the reaction of SnCl
2
with elemental sulfur at 170 °C in oleylamine.
Scanning electron microscopy (SEM) images showed tetrahedral nano- and microcrystals
ranging between 200 and 300 nm on each side (Fig. 1.3.a). Modifications in the reaction
time or amine surfactant (decyl-, dodecyl-, hexadecyl-, and oleylamine) did not have an
effect on the size, polydispersity, or shape of the SnS product. The metastable zinc
blende crystal structure (a = 5.85 Å) was verified by powder X-ray diffraction (XRD);
however, small amounts of Sn, sulfur, Sn(OH)
2
and orthorhombic SnS were also
observed (Fig. 1.3.b). Up to this point, the zinc blende phase had only been observed in
bulk SnS grown epitaxially on a NaCl seed layer and it had been thought to be mixed
with the orthorhombic α- SnS phase.
78,79
Selected-area electron diffraction (SAED)
confirmed that the SnS nano- and microcrystals were single crystalline; while a lattice
spacing of d = 8.1 ± 0.2 Å was observed for the (110) lattice fringes by high-resolution
transmission electron microscopy (TEM). Interestingly the Schaak group studied the
morphology-dependent polymorphism of SnS using an in-depth crystallographic analysis
that correlates high-resolution TEM data of individual nanocrystals with ensemble-based
electron diffraction and powder XRD data and found that the XRD peaks which are often
interpreted as a mixture of metastable zincblende-type SnS and α-SnS are better
described by a single phase with pseudotetragonal modification.
56
Greyson et al. also
explored the thermal stability of their as-synthesized SnS by heating the nano- and
microcrystals under Ar at 300 °C for 3 h, and observed no change in the SnS morphology
!
!
12
or crystal structure. In contrast, when the SnS nano- and microcrystals were heated in the
presence of oleylamine at 250 °C for 3 h, a nearly complete conversion (>90%) from zinc
blende to orthorhombic platelets was observed. This phase transformation in solution
may be a result of the oleylamine lowering the activation barrier for structural
rearrangement.
The orthorhombic SnS platelets possessed a similar absorption profile to that of bulk
orthorhombic SnS, with a strong absorption onset at ca. 980 nm (corresponding to the
direct band gap at E
g
= 1.3 eV) and a weaker absorption edge near 1100 nm
(corresponding to the indirect band gap at E
g
= 1.1 eV). The calculated band structure of
zinc blende SnS predicts that the material is either a metal or a small indirect band gap
semiconductor; however, the absorption profile of the zinc blende tetrahedra is blue
shifted to 700 nm (E
g
= 1.8 eV) relative to the orthorhombic phase. The higher energy
band gap for the zinc blende phase relative to the orthorhombic phase may be a result of
the different symmetries of the crystal structures; however, the impurities observed by
XRD in the zinc blende SnS may also play a role in the spectral blue shift.
In 2008, Hickey et al. synthesized monodisperse and slightly sulfur-rich sub-10 nm SnS
nanocrystals with a narrow 10% size distribution (Fig. 1.4.).
80
The SnS nanocrystals were
synthesized by the hot injection of thioacetamide in oleylamine into a mixture of
Sn[N(SiMe
3
)
2
]
2
, oleic acid, trioctylphosphine, and octadecene at 170 °C. Variation of the
oleic acid/oleylamine ratio allowed for shape control. A 1 : 2 ratio of oleic
acid/oleylamine produced spherical nanocrystals, whereas a 1 : 1 ratio produced faceted
!
!
13
nanocrystals with a distinct triangular or truncated triangular projection. Between these
two extremes, the nanocrystals were found to be monodisperse, but their morphology was
difficult to define. It was also found that as the concentration of oleic acid increased, the
nanocrystals became larger in size. The authors determined that other Sn
2+
precursors,
such as Sn(OAc)
2
and Sn(oleate)
2
, produced material with a large sheet-like morphology
rather than small monodisperse nanocrystals.

Figure 1.4. (a) TEM images of spherical SnS nanocrystals synthesized from Sn[N(SiMe
3
)
2
]
2

and
thioacetamide. (b) High-resolution TEM image of a single SnS nanocrystal. Reprinted with
permission from ref. 80 (Copyright 2008 American Chemical Society).

The nanocrystals were found to possess the orthorhombic α-SnS crystal structure (a =
4.31 Å, b = 11.26 Å, c = 3.98 Å) typical of these materials; however, a small number of
peaks attributed to the zinc blende phase and surface-oxidized SnO were also observed
by XRD. The absorption profile of clear solutions of the spherical SnS nanocrystals was
very steep and blue shifted as compared to bulk SnS, with a calculated indirect band gap
of E
g
= 1.6 eV. While the bulk material possesses both direct and indirect band gaps that
are relatively close in energy (i.e., 200–600 meV), the authors suggest that the positions
!
!
14
of these two band gaps relative to one another may change as a function of quantum
confinement. The applicability of these SnS nanocrystals as a photoconductive material
was assessed through a photoelectrochemical experiment whereby a thin film of the
nanocrystals was irradiated with 470 nm light and the photocurrent was measured (Fig.
1.5.). It was found that the photocurrent response for the SnS nanocrystals was stable and
repeatable over many light/dark cycles.

Figure 1.5. Photocurrent response of a thin film of SnS nanocrystals on an ITO substrate in 0.1
M sodium sulfate electrolyte (E
app
= 800 mV vs. Ag/AgCl). Reprinted with permission from ref.
80 (Copyright 2008 American Chemical Society).

In 2009, Xu et al. synthesized sub-5 nm SnS nanocrystals that were dispersible in polar
solvents, such as ethanol.
81
In this method, SnBr
2
was reacted with Na
2
S in ethylene
glycol at room temperature in the presence of various stabilizing ethanolamines as
ligands. Three different ethanolamines were used as stabilizing ligands: triethanolamine
(TEA), N-methyldiethanolamine (MDEA), or N,N-dimethylethanolamine (DMEA).
Among the three ethanolamines surveyed, TEA (with three hydroxyl groups) produced
!
!
15
the smallest and most monodisperse SnS nanocrystals. It is thought that TEA plays dual
roles of (1) coordinating to the Sn
2+
to form a [Sn(TEA)n]
2+
precursor complex, and (2)
strongly binding to the SnS nanocrystal surface upon nucleation through the multiple
hydroxyl groups. As the number of hydroxyl groups decrease in the ethanolamines, the
resulting nanocrystals gradually get larger and less monodisperse with TEA yielding
nanocrystals of 3.2 ± 0.5 nm, MDEA yielding nanocrystals of 4.0 ± 2.0 nm, and DMEA
yielding nanocrystals of 5.0 ± 4.0 nm in diameter.
The highly crystalline nature of the SnS nanocrystals was demonstrated by observation of
atomic lattice fringes by HR-TEM for apparent single-crystalline particles. SAED was
used to confirm that the nanocrystals were in the expected orthorhombic α-SnS phase (a
= 11.14 Å, b = 3.97 Å, and c = 4.34 Å), with no other phases being observed.
Transmission spectra on colloidal 5.0 nm SnS nanocrystal dispersions indicated an
indirect band gap at E
g
= 1.1 eV that is similar to the bulk, and no fluorescence was
observed for these SnS nanocrystals, which may be consistent with their indirect band
gap behavior.
In 2010, Liu et al. synthesized SnS nanocrystals that were single crystalline,
monodisperse, and size-tunable (Fig. 1.6.).
82
The nanocrystals were synthesized by the
hot injection of S(SiMe
3
)
2
in octadecene into a solution of SnCl
2
in oleylamine at 200 °C.
After the addition of oleic acid, the SnS nanocrystals remained dispersed in toluene for
more than 6 months if stored under an inert atmosphere. Nanocrystal size can be
controlled between 6, 12, and 20 nm by varying injection and growth temperatures
!
!
16
between 120, 150, and 210 °C, respectively. It was also determined that a 1 h incubation
yielded monodisperse SnS nanocrystals, while a 5 min incubation yielded a larger size
distribution suggesting a size-focusing growth mechanism.

Figure 1.6. SEM image of 6-nm SnS nanocrystals synthesized from SnCl
2
and S(SiMe
3
)
2
.
Reprinted with permission from ref. 82 (Copyright 2010 IOP Publishing).

The resulting nanocrystals were found to be in the orthorhombic β-SnS phase (a = 4.14
Å, b = 11.49 Å, c = 4.17 Å), as determined by XRD analysis. No band gap shift was
observed between the three differently sized samples (6, 12, and 20 nm), which all
showed an absorption onset at ca. 1.55 eV similar to that reported by Hickey et al.;
80

however, Liu et al. assigned it to a direct band gap transition rather than an indirect band
gap transition.
Most recently, Ning et al. reported the synthesis of SnS nanostructures with various sizes
and morphologies.
83
In this preparation, Sn
6
O
4
(OH)
4
was used as the Sn
2+
precursor. The
Sn
6
O
4
(OH)
4
precursor was dissolved in oleic acid and oleylamine, and then
thioacetamide was injected with oleylamine at elevated temperatures ranging from 120–
!
!
17
150 °C. If the thioacetamide was injected at 150 °C with a 1 : 1 molar ratio of Sn/S, then
5 nm SnS nanocrystals were produced. If the thioacetamide was injected at 120 °C with
a 2 : 1 molar ratio of Sn/S, then 13 nm, single-crystalline SnS nanoflowers were
produced through oriented attachment. As the reaction time was increased from 3–10
min, the SnS nanoflowers converted to amorphous SnS with a sheet-like morphology.
The authors do not discuss whether the nanocrystals are in the expected 1 : 1
stoichiometry for SnS, or if varying the Sn/S ratio makes the resulting nanocrystals either
sulfur- or tin-rich.
XRD analysis of the nanocrystals and nanoflowers confirmed they were in the expected
orthorhombic α-SnS crystal structure (a = 4.33 Å, b = 11.19 Å, and c = 3.98 Å).
Absorption spectra of the 5 nm SnS nanocrystals suggest both indirect and direct optical
band gaps of E
g
= 1.6 and 3.6 eV, respectively. Emission spectra produced by excitation
between 340 and 380 nm appear to confirm the direct band gap in SnS with an emission
maximum at 3.2–3.6 eV. Both the indirect and direct band gaps are blue shifted as a
result of quantum confinement effects; however, the direct wide band gap is blue shifted
substantially more than that of the indirect band gap from bulk values.
1.4.2. Tin Selenide Nanocrystals
Similar to SnS, SnSe has also proven to be a promising lead-free IV–VI PV material.83
Bulk SnSe exhibits p-type conductivity, with hole mobilities up to 10
3
cm
2
V
−1
s
−1
along
the c axis at −196 °C.
85,86
The indirect and direct band gaps (Eg = 0.90 eV and 1.30 eV,
!
!
18
respectively) of bulk SnSe correlate well with the optimum band gap values for solar
cells, which fall between Eg = 1.0 and 1.5 eV.
49
In 2007, Pejova and Grozdanov
demonstrated the effects of quantum confinement in SnSe by studying the optical
properties of SnSe nanocrystals deposited as a thin film with an average grain size of 14.8
nm.
86
The as-deposited nanocrystals possessed indirect and direct band gaps of Eg = 1.20
and 1.74 eV, respectively, which are both blue shifted from the band gaps of the bulk
material. This suggests that the nanocrystalline grains were smaller than the Bohr exciton
radius for SnSe. Upon annealing the films to 150 °C for 1 h, the average grain size
increased to 23.3 nm, accompanied by a red shift of the indirect and direct band gaps to
E
g
= 1.10 and 1.65 eV, respectively. Promptly after demonstrating quantum confinement
effects in nanocrystalline SnSe thin films, Pejova and Tanusevski went on to demonstrate
the potential of SnSe as an active component in PVs by studying the charge-transport
properties of these films.
88
Their findings showed that SnSe is a photoconductive material
with contributions from both indirect and direct band transitions that are close in energy,
which agrees with previous work.
49
The dominant charge carriers are holes, indicating
that the SnSe semiconductor films are p-type and act as acceptors when in a solar cell
device. Thermionic emission over the crystal grain boundaries was determined to be the
predominant charge transport mechanism at room temperature. The average lifetime of
the minority charge carriers (electrons) was relatively high (1.78 ms), giving further
favorable evidence for using SnSe thin films as absorber materials in solar cells.
!
!
19
A number of solution-phase synthetic routes to SnSe nanocrystals have been reported
over the past decade; however, many of these reports lacked the synthetic control needed
to produce well-defined nanocrystals.
89-92
In 1999, Wang et al. reported a mild, low
temperature reductive route to ill-defined and morphologically diverse SnSe nanorods
(approximate dimensions 30 nm × 1.5 µm) using an ethylenediamine chelate to direct
growth.
89
In 2000, Zhang et al. published an aqueous route to nanocrystalline SnSe with
large, sheet-like morphologies.
90
The product precipitated immediately upon the reaction
of a highly alkaline aqueous solution of selenium and a mixture of SnCl
2
with tartaric
acid. It is possible that the absence of structure-directing ligands other than the tartrate
present in the alkaline mixture led to the large, ill-defined nanocrystalline product. In
2003, Shen et al. synthesized the first SnSe nanowires with high aspect ratios (i.e.,
150) and a narrow size distribution.
91
They used a simple, rapid ethylenediamine-assisted
polyol process at 200 °C with selenium and SnCl
2
. They believe the ethylenediamine
was the key factor to obtain phase pure SnSe nanowires because it reduces the elemental
selenium to form Se
2−
. The morphology of the SnSe product is also highly dependent on
the presence of ethylenediamine. Without the addition of ethylenediamine, the SnSe
product was found to contain unreacted elemental selenium and showed flake and particle
morphology. They believe ethylenediamine forms a precursor [Sn(en)
2
]
2+
complex that
serves as a molecular template whereby selenium ions may coordinate to form one-
dimensional SnSe nanorod structures; however, a high degree of particulate product is
still observed with the introduction of ethylenediamine. In 2004, Han et al. reported the
room temperature preparation of SnSe nanorods in a similar route to Zhang et al.
90,92

!
!
20
Nanorods of SnSe were produced by mixing a highly alkaline aqueous selenium solution
with SnCl
2
in the presence of trisodium citrate. The agglomerated nanorods had an
average diameter of 90 nm and lengths up to 1 µm. As with the Zhang synthesis, the
highly alkaline conditions (>10 M NaOH) needed to completely dissolve the Se make the
synthesis method less than ideal.

Figure 1.7. TEM images of SnSe nanocrystals synthesized from SnCl
2
and
t
Bu
2
Se
2
. (a)
High- resolution TEM image of a single nanocrystal. (b) SAED pattern for an ensemble
of SnSe nanocrystals. (c) Low-resolution TEM image of SnSe nanocrystals. Reprinted
with permission from ref. 94 (Copyright 2010 American Chemical Society).

In 2002, Schlecht et al. took a different approach to the solution-phase SnSe synthesis by
employing diorganodichalcogenides as soluble chalcogenide sources.
93
They turned to
diphenyl diselenide (Ph
2
Se
2
) because of its solubility in diglyme, after they were unable
to obtain nanocrystalline SnSe by directly reacting Sn
0
with elemental selenium. The
!
!
21
overall reaction required two-steps. In the first step, 2 equiv of Ph
2
Se
2
reacted with
activated tin at 65 °C in THF to produce a Sn(SePh)
4
selenolate rather than the intended
SnSe. In the second step, thermolysis of Sn(SePh)
4
at 300 °C leads to the formation of
nanocrystalline SnSe with byproducts of Ph
2
Se and Ph
2
Se
2
. The resulting nanocrystals
possessed a broad size distribution (3–50 nm) with no control over particle morphology.
Recent advancements in the solution-phase synthesis of well-defined SnSe nanocrystals
were made by Franzman et al. in 2010 with the first publication of small colloidal SnSe
nanocrystals shown to exhibit quantum confinement effects.
94
Their synthetic route
involved the use of a diorganodichalcogenide as the chalcogen source.
95-100
A
stoichiometric amount of di-tert-butyl diselenide (
t
Bu
2
Se
2
) was injected into a solution of
anhydrous SnCl
2
, dodecylamine, and dodecanethiol at 95 °C. Following injection, the
reaction temperature was raised to 180 °C for 4 min and then quenched by cooling to
obtain phase-pure SnSe nanocrystals. They found that control over the nanocrystal
composition, and more specifically the oxidation state of tin, could easily be obtained by
controlling the amount of
t
Bu
2
Se
2
added to the reaction. Addition of 0.5 equiv of the
t
Bu
2
Se
2
gives phase pure SnSe, while addition of 1.0 equiv of the
t
Bu
2
Se
2
gives SnSe
2
in
a result similar to Schlecht et al.
93

Transmission electron microscopy analysis revealed the SnSe product to be composed of
elongated anisotropic nanocrystals of variable length and consistent width (19.0 nm ± 5.1
nm; Fig. 1.7.). The product was phase pure and crystallized in the typical orthorhombic
phase (a = 11.55 Å, b = 4.16 Å, c = 4.45 Å) with a distorted rock salt structure. A 48 : 52
!
!
22
tin to selenium ratio, with Sn
2+
and Se
2−
oxidation states, was confirmed through a
combination of energy-dispersive X-ray spectroscopy (EDX) and X-ray photoelectron
spectroscopy (XPS). The SnSe nanocrystals absorbed through the visible spectrum and
into the near-IR having a direct band gap (E
g
= 1.71 eV) which was blue-shifted relative
to the bulk (E
g
= 1.30 eV) due to quantum confinement effects. Given the potential of
these quantum confined SnSe nanocrystals, their utility in a hybrid PV device was
explored (vide infra).
The size, shape, and surface chemistry of quantum confined SnSe nanocrystals were
further investigated by Baumgardner et al.
101
They successfully carried out a solution-
phase synthesis of SnSe through hot (65–175 °C) injection of Sn[N(SiMe
3
)
2
]
2
into
TOPSe in the presence of oleylamine. After nucleation, oleic acid was introduced to the
mixture and then the reaction was quenched. The resulting SnSe nanocrystals were
unagglomerated and quasispherical in shape, as revealed by TEM analysis. It was
observed that the SnSe nanocrystal shape is sensitive to the surface ligands present.
When oleic acid was present prior to precursor injection, nucleation was inhibited. This
finding differs from previous PbSe/Te syntheses, where oleic acid was used to tune the
nucleation step.
20,102
For the SnSe synthesis, Baumgardner et al. attributed the inhibition
of nucleation with oleic acid to the high binding affinity of oleate for Sn
2+
, and the
resulting lowered chemical potential driving force for nucleation. When oleic acid was
injected after nucleation, it was found that growth was accelerated until the equilibrium
!
!
23
size was reached. Replacing oleic acid with dodecanethiol resulted in the growth of
anisotropic SnSe nanocrystals, similar to those produced by Franzman et al.
94

Figure 1.8. (a) Low-resolution TEM image of SnSe nanocrystals synthesized from
Sn[N(SiMe
3
)
2
]
2
and TOPSe. (b) High-resolution TEM image of a single nanocrystal. (c) The
(001) and (010) projections of a SnSe unit cell with Pnma symmetry. Reprinted with permission
from ref. 101 (Copyright 2010 American Chemical Society).

With this synthetic method, SnSe nanocrystals from 4 to 10 nm could be controllably
synthesized by tuning the injection and reaction temperatures (Fig. 1.8.), and
manipulation of the reaction temperature was also found to affect the crystal structure.
Reaction temperatures held at 175 °C produced α-SnSe nanocrystals exhibiting the Pnma
crystal structure; however, when the reaction temperature was lowered to 105 °C, a
decrease of the nanocrystal diameter was observed with a concomitant increase in the d-
spacings of several reflections. The observed increase in d-spacing suggests that the SnSe
crystal structure changed from Pnma to Cmcm symmetry, which is typically only
!
!
24
observed after high temperature annealing of bulk SnSe to 600 °C. The formation of β-
SnSe nanocrystals with the Cmcm crystal structure is further supported by the
disappearance of the characteristic Pnma (020) and (112) reflections in the XRD pattern.
This is an interesting result because a metastable phase is being observed at lower
temperatures.
54,57,103,104

Figure 1.9. (a) Tauc plot from UV-vis-NIR data for SnSe nanocrystals showing an indirect band
gap. (b) Approximate 1r
-2
relationship between band gap and mean nanocrystal size. Reprinted
with permission from ref. 101 (Copyright 2010 American Chemical Society).

Nanocrystals in the size range from 4–9 nm exhibited quantum confinement effects,
similar to the results observed by Franzman et al. The indirect band gap varied from E
g
=
1.2–0.9 eV for nanocrystal sizes ranging from 4–9 nm (Fig. 1.9.a), respectively, while the
direct band gap varied from E
g
= 1.8–1.3 eV over the same size range. Both the indirect
!
!
25
and direct band gaps demonstrated a rough 1/r
2
dependence on the nanocrystal size (Fig.
1.9.b). Although a prototype solar cell was not reported, the Hanrath group did report the
applicability of their SnSe nanocrystals as a photoconductive material. The
photoconductivity of formic acid passivated SnSe films over interdigitated gold
electrodes was confirmed through transient current–voltage (I–V) characteristics.
Transient photocurrent was observed at a bias of 2 V µm
−1
under 100 mW cm
−2

illumination; however, significant signal degradation occurred over time. This
photocurrent degradation was attributed to the photooxidation of organic species (i.e.,
formic acid or oleic acid) bound to the nanocrystal surface.
The synthesis of colloidal single-crystal nanosheets has attracted much interest because
of the inherent anisotropy exhibited by two-dimensional structures and the ensued wide
range of potential applications that high quality nanosheets possess. Schaak et al.
demonstrated control over shape, lateral dimensions, and sheet thickness by reacting
SnCl2, trioctylphosphine selenide (TOP-Se), and hexamethyldisilazane (HMDS) in
oleylamine using a one-pot approach.
105
After slowly heating the mixture to 240 °C and
holding this temperature for 30 min, they obtained orthorhombic α-SnSe nanostructures
that exhibited a uniform square-like morphology (500 nm × 500 nm). Control over
nanosheet thickness between 10 and 40 nm while maintaining uniform lateral dimensions
was achieved by tuning the concentrations of SnCl
2
and TOP-Se. The resulting material’s
band gap of Eg = 1 eV and preliminary photoresponse measurements were determined
from SnSe nanosheet samples that were drop casted.
!
!
26
1.4.3. Tin Telluride Nanocrystals
Narrow band gap IV–VI semiconductor nanocrystals can be used for NIR-absorbing PV
devices (e.g., in a tandem cell); however, it has been a challenge to synthesize
nanocrystals with band gap energies below 0.5 eV.
106,107
Bulk SnTe is isotropic, which
allows for relatively high p-type conductivity with hole mobilities of 840 cm
2
V
−1
s
−1
at
27 °C.
108
Tin telluride is a direct band gap semiconductor that exhibits a narrow band
gap of E
g
= 0.2 eV at room temperature,
109
and the observation of quantum-size effects in
SnTe nanocrystals by Kovalenko et al. has sparked interest in this material as a stable and
less toxic semiconductor in PV applications.
110

One of the first solvothermal syntheses of tin monochalcogenides entailed the use of
soluble diaryl dichalcogenides and activated Sn
0
nanocrystals.
93
This work evolved from
the use of elemental selenium and tellurium to the use of diphenyl diselenide (Ph
2
Se
2
)
and ditelluride (Ph
2
Te
2
) to obtain single-phase SnSe and SnTe micro- and nanocrystals.
All syntheses began with the known reduction method of SnCl
2
by Li[Et
3
BH] in THF and
diglyme to produce the Sn
0
nanocrystal precursor.
111
The reaction of Sn
0
with Ph
2
Te
2

produced two different morphologies of the crystalline material. Concentrated reaction
mixtures produced large 60 nm agglomerates with random orientation; but as the reaction
concentration decreased, smaller 15 × 40 nm star-shaped nanocrystals were formed. The
cubic rock-salt structure (a = 6.33 Å) of the SnTe nanocrystals was determined by
HRTEM, SAED and XRD to closely match that of the bulk SnTe, and the 1 : 1
stoichiometry of Sn : Te was confirmed by EDX. As discussed previously, the equivalent
!
!
27
reaction with Ph
2
Se
2
yielded the discrete selenolate Sn(SePh)
4
, and a short pyrolysis at
300 °C was needed to convert the product into SnSe nanocrystals.

Figure 1.10. (a,b) TEM images of 10-nm SnTe nanocrystals synthesized from Sn[N(SiMe
3
)
2
]
2
and TOPTe. Reprinted with permission from ref. 110 (Copyright 2007 American Chemical
Society).

In 2007, Kovalenko et al. synthesized monodisperse SnTe nanocrystals in solution that
were tunable in size between 4.5 and 15 nm (Fig. 1.10.), with their corresponding band
gaps ranging between E
g
= 0.8 and 0.38 eV, respectively.
110
The stoichiometric SnTe
nanocrystals were prepared by the reaction of Sn[N(SiMe
3
)
2
]
2
in octadecene with TOPTe
in oleylamine at 150 °C, with oleic acid being added to passivate the surface of the
resulting nanocrystals. The uniform and nearly spherical shape of the nanocrystals is
evident from the TEM analysis, while the cubic rock-salt crystal structure (a = 6.24 Å)
was confirmed by XRD and HRTEM. Increasing the concentration of oleylamine and the
temperature of injection/incubation generally resulted in larger SnTe nanocrystals.
The absorption spectra of the SnTe nanocrystals were studied to determine the size effect
on the optical band gap. The resulting nanocrystals possessed broad excitonic peaks in
!
!
28
the IR region that blue-shifted with decreasing size, consistent with quantum confinement
effects. The optical band gaps for the 14 and 7.2 nm SnTe nanocrystals were E
g
= 0.54
and 0.39 eV, respectively; these values are near the calculated optimal value of 0.35 eV
for semiconductors that may exhibit MEG.
112
The charge transport properties of thin
films of the SnTe nanocrystals were studied to evaluate their potential in PV and
thermoelectric applications. The resulting low electrical conductivities of σ ≈ 10
−10
S
cm
−1
were likely the result of the large interparticle spacing in the films, but the
conductivity increased by almost six orders of magnitude when the films were treated
with hydrazine in anhydrous acetonitrile (Fig. 1.11.). The hydrazine treated films showed
n-type conductivity, indicated by the increase in conductivity when a positive bias was
applied to the back gate electrode.
More recently, Ning et al. were able to synthesize SnTe nanocrystals and nanowires.
113

The unique properties exhibited by nanorods and nanowires make these low dimensional
materials attractive,
114
and have motivated studies on their formation via aggregation of
0-D nanocrystals by oriented attachment.
115
The synthesis of the SnTe nanocrystals was
achieved by the hot injection of TOPTe into a mixture of Sn
6
O
4
(OH)
4
, oleic acid, and
oleylamine (or octylamine) at 180 °C. The product was incubated at 165 °C, and formed
stable concentrated colloidal solutions upon purification. Transmission electron
microscopy images confirm that when oleylamine is used, the resulting SnTe
nanocrystals were ca. 4 nm in diameter. When the oleylamine was replaced by shorter
octylamine ligands, the resulting nanocrystals were larger ( 8 nm) in size and of low
!
!
29
crystallinity; however, these nanocrystals transformed into 50 nm long crystalline
nanowires at extended reaction times through oriented attachment. The authors speculate
that the shorter chain amine induced a faster growth rate of the nanocrystals that in turn
caused a larger size and lower crystallinity when compared to the longer chain
oleylamine.
116
The crystalline SnTe nanowires grew along the (100) direction, and at
extended reaction times, the SnTe nanowires reached 150 nm in length and 10 nm in
width.

Figure 1.11. I-V curves for films of the (a) oleic acid capped (b) hydrazine treated SnTe
nanocrystals. Upon treatment with hydrazine, the SnTe nanocrystals demonstate n-type behavior.
Reprinted with permission from ref. 108 (Copyright 2007 American Chemical Society).

!
!
30
Reiss et al. demonstrated that charge transfer and charge transport processes strongly
depend on nanocrystals’ surface quality and that prior to the application of tin
chalcogenide nanocrystals in optoelectronics further surface engineering is required.
121

The group used 119Sn-Mössbauer spectroscopy to show that, regardless of the tin or
sulfur reagents used, the Sn(IV) to Sn(II) ratio in SnS nanocrystals changed from 20:80 to
40:60 after five minutes of air exposure. Similarly, the Sn(IV) to Sn(II) ratio in SnSe
nanocrystals changed to 43:57 after air exposure; while a higher tendency to oxidize was
found in SnTe nanocrystals that were exposed to air, as the Sn(IV) to Sn(II) ratio changed
to 55:45.

1.4.4. Germanium Sulfide and Selenide Nanocrystals
Similar to the tin analogs, both bulk GeS and GeSe are native p-type semiconductors; for
example, hole mobilities of 90 cm
2
V
−1
s
−1
have been measured along the c axis for GeS at
27 °C.
118,119
Vaughn et al. have recently reported the first solution chemistry route for
colloidal GeS and GeSe nanostructures.
120
The complicated band structure and the
closeness of the direct and indirect band gaps of GeS and GeSe result in a range of values
for both the direct and indirect band gaps (Eg = 1.6–1.7 eV for GeS and Eg = 1.1–1.2 eV
for GeSe) that overlap well with the solar spectrum.
66,67
The GeS nanosheets were
synthesized via the reduction of GeI
4
in hexamethyldisilazane, oleylamine, oleic acid, and
dodecanethiol at 320 °C for 24 h. To synthesize GeSe nanosheets, TOPSe was used as
the selenium precursor in place of dodecanethiol. TEM images show mainly elongated
!
!
31
hexagons (2–4 µm by 0.5–1 µm; Fig. 1.12.), with thicknesses between 3 and 20 nm for
GeS and 5–100 nm for GeSe (as estimated by AFM). The phase-pure orthorhombic GeS
and GeSe nanosheets are both oriented along the [100] direction. The lattice parameters
for GeS (a = 10.52 Å, b = 3.65 Å, and c = 4.30 Å) and GeSe (a = 10.78 Å, b = 3.81 Å,
and c = 4.37 Å) corresponded well with literature values.

Figure 1.12. (a) TEM images of GeS and GeSe (inset) nanosheets synthesized from GeI4 and
dodecanthiol or TOPSe, respectively. SAED patterns for (b) an ensemble of GeS nanosheets and
(c) a single crystalline GeS nanosheet oriented along the [100] direction. Reprinted with
permission from ref. 120 (Copyright 2010 American Chemical Society).

Diffuse reflectance spectroscopy was used to approximate the indirect and direct band
gaps of GeS (E
g
= 1.58 and 1.61 eV, respectively) and GeSe (E
g
= 1.14 and 1.21 eV,
!
!
32
respectively), which are very close to the values of the bulk band gap. Four-point I–V
measurements were used to measure a conductivity of σ = 4.7 × 10
−6
S cm
−1
for drop-cast
thin films of the GeSe nanosheets, while a two-point I–V measurements demonstrate the
p-type character of this material (Fig. 1.13.). The conductivity value is comparable to
other colloidal nanocrystalline semiconductor thin films.
121

1.5. Photovoltaic Device Applications
Recent advancements in the development of simple, reproducible and low-cost synthetic
techniques towards high-quality tin and germanium monochalcogenide nanocrystals have
led to the end goal of their inclusion into thin film PV device architectures. To date, the
reported devices remain limited to the inclusion of either SnS or SnSe nanocrystals into
various device architectures; however, the data appear promising thus far. For example,
Stavrinadis et al. showed that inclusion of SnS nanocrystals into a lead chalcogenide type
II bilayer heterojunction solar cell resulted in larger open-circuit voltages (V
OC
= 0.44 V)
as compared to Schottky cells based on pure PbSe or PbS nanocrystals.
122
In this work,
SnS nanocrystals were incorporated into an ITO/SnS/PbS/Al device stack. The open-
circuit voltage was >100% larger than that of the control Schottky ITO/PbS/Al device
(V
OC
= 0.20 V). They attributed the increase in open circuit voltage of the device to the
built-in electric field of the SnS/PbS heterojunction, which acts as an electron blocking
layer that assists in the diffusion of charge carriers to their respective contacts. The short-
circuit current, fill factor and overall power conversion efficiencies (J
SC
= 1.84 mA cm
−2
,
!
!
33
FF = 0.30, η
P
= 0.31%) of the bilayer device were generally found to be lower than other
reported lead chalcogenide Schottky devices.
34
They believe the key to future
optimization (i.e., increasing J
SC
) lies in improving the synthesis and post-synthesis
processing techniques of the nanocrystals; however, their results do suggest that SnS
nanocrystalline films can lead to substantially improved properties of multilayer PVs.
Wang et al. showed that the addition of SnS nanocrystals into hybrid polymer containing
bulk heterojunction PV cells improves the device performance when compared to pristine
polymer devices lacking the SnS.
123
Their device was structured as
ITO/PEDOT:PSS/SnS:polymer/Al, where the active layer was synthesized by blending
the nanocrystals with poly[2-methoxy-5-(3′,7′-dimethyloctyloxy)-1,4-phenylenevinylene]
(MDMO-PPV) or poly(3-hexylthiophene) (P3HT). Both active layers gave
improvements in the absorption intensity and range of the absorption spectrum upon
incorporation of the SnS nanocrystals. The power conversion efficiency (η
P
= 1.08 ×
10
−2
%) of the SnS:P3HT solar cells was highest at 86 wt% SnS, due to an increased
short-circuit density (J
SC
= 0.026 mA cm
−2
vs. 0.017 mA cm
−2
for neat polymer);
however, this came at the cost of V
OC
and fill factor due to changes in P3HT morphology
upon addition of SnS. The performance of the SnS:MDMO-PPV active layer showed
greater improvements in power conversion efficiency upon addition of SnS nanocrystals,
specifically at 67 wt% SnS (η
P
= 2.05 × 10
−2
%), without the decreases in V
OC
and fill
factor observed for P3HT cells. At 67 wt% SnS in MDMO-PPV, the power conversion
efficiency was 26.6 times that of neat polymer cell.
!
!
34
More recently, Wang et al. showed improvement by one order of magnitude on the power
conversion efficiency of their previously reported SnS:MDMO-PPV PV cell by using
SnS/SnO heterojunction nanocrystals rather than pure SnS nanocrystals.
124
The
improved performance of this device is a result of the unique SnS/SnO rod-like
morphology (as compared to the 0-D SnS morphology), which allows for more facile
charge transport within the inorganic phase and improved current density in the device.
Moreover, the addition of SnO into the inorganic phase widens the band gap and, as a
result, V
OC
increases from 0.37 V for the SnS:MDMO-PPV device to 0.74 V for the
SnS/SnO:MDMO-PPV device.

Figure 1.13. (a) Four-point I-V curve giving a conductivity of σ = 4.7x10
-6
S cm
-1
for dropcast
films of the GeSe nanosheets. (b) Two-point I-V curve giving turn-on potentials of -6.5 and +10
V suggesting p-type character for the GeSe nanosheets. Reprinted with permission from ref. 120
(Copyright 2010 American Chemical Society).

The first demonstration of PV application of a cell based on TiO
2
/SnS films was reported
in 2010 by Wang et al.
125
An electrochemical solar cell structure of FTO/Pt + electrolyte
+ SnS/TiO
2
/FTO yielded a high V
OC
= 0.471 V, a JSC = 0.30 mA cm
−2
, a power
!
!
35
conversion efficiency of η
P
= 0.10%, and FF = 0.71 under 1 sun illumination. Power
conversion efficiencies dropped to η
P
= 0.03% without the presence of the
nanocrystalline SnS layer, proving that addition of the SnS results in better PV
performance. The authors mention that this SnS/TiO
2
device compares favorably with
previously published CIS/In
2
S
3
/TiO
2
device architectures.
126

Tin selenide nanocrystals have also recently been used as PV materials in hybrid
polymer/nanocrystal device architectures. In 2010, Franzman et al. demonstrated the
utility of SnSe nanocrystals by integrating them into a conducting polymer as the electron
accepting layer in a hybrid PV device.
94
Their device consisted of a SnSe:poly[2-
methoxy-5-(3′,7′-dimethyloctyloxy)-1,4-phenylenevinylene] (MDMO-PPV) absorbing
layer, a perylene-3,4,9,10-tetracarboxylic diimide (PTCDI) acceptor/hole-blocking layer,
and a LiF/Al bilayer cathode on a glass substrate. When compared to analogous neat
polymer devices, it was found that the J
SC
of the SnSe hybrid cell was nearly twice that of
the neat MDMO-PPV cell. Both the V
OC
and fill factors of the two were comparable
suggesting the absence of charge trapping on the SnSe nanocrystals. The power
conversion efficiency was improved by 100% (η
P
= 0.03% for neat MDMO-PPV
compared to η
P
= 0.06% for SnSe/MDMO-PPV) upon inclusion of SnSe into the polymer
(0.25 : 1.0, w/w, SnSe/MDMO-PPV). The external quantum efficiency doubled near
500 nm with the absorption coefficient remaining nearly the same at that wavelength for
the hybrid device compared to the neat polymer. These data indicate that the polymer
acts as the primary absorber and the observed increase in power conversion efficiency is
!
!
36
likely a result of electron transfer from MDMO-PPV to the SnSe nanocrystals. It should
be noted that this was the first reported synthesis and utilization of quantum confined
SnSe nanocrystals in a PV device, and that much improvement is still needed in order for
this material to function as an earth abundant absorber material.
1.6. Conclusions
A variety of tin and germanium monochalcogenide nanocrystals have been synthesized
over the past decade. The majority of these syntheses utilize a hot-injection type approach
whereby a chalcogenide source (e.g., trioctylphosphine chalcogenide,
diorganodichalcogenide, thio/selenocarbonyl, dissolved elemental source, etc.) is injected
into a hot solution of the metal precursor (e.g., metal salt, metal amide, dissolved metal
oxo cluster, etc.) and various stabilizing ligands. These methods have yielded
nanocrystals of various sizes and shapes, with varying degrees of quality in terms of
monodispersity, morphological fidelity, and phase purity. The first major challenge in
this chemistry results from the tendency of the GeS, GeSe, SnS and SnSe nanocrystals to
form sheet-like structures (similar to those synthesized by Vaughn et al.) instead of 0-D
particles. This morphological effect is a direct result of the Pnma double layer crystal
structure that these materials adopt, as opposed to the rock salt structure adopted by SnTe
and the lead chalcogenides. The second major challenge in this chemistry is the synthesis
of phase pure nanocrystals, since metastable phases can be easily accessed, in addition to
the crystalline impurities that result from oxidation of tin and germanium to the Sn
4+
and
Ge
4+
oxidation states. While a great deal of progress has been achieved in overcoming
!
!
37
these challenges, the same level of size control and monodispersity that has been realized
in the lead chalcogenides has still not been achieved in these systems. Moreover, there
has not been much success in controlling the dimensionality of a given tin or germanium
monochalcogenide nanocrystal between 0-D, 1-D, 2-D and 3-D type structures.
While there have been some notable reports of using SnS in heterojunction solar cells
(e.g., η
P
= 1.3% for a SnO
2
/SnS/CdS/In device structure69), the majority of PV devices
derived from SnS or SnSe nanocrystals have thus far demonstrated rather poor power
conversion efficiencies (η
P
< 1%). The single greatest challenge in increasing the
performance of these PV devices is to achieve a greater degree of control over the SnS
and SnSe nanocrystal surface. Significant device improvements will be made by
removing the insulating ligands protecting the nanocrystals and replacing them with
small molecules, such as hydrazine or metal chalcogenide clusters.
127,128
This will allow
for much improved interparticle coupling (i.e., charge transport through the
polycrystalline layer), which is a function of interparticle spacing and nanocrystal surface
chemistry.
121
Moreover, if MEG is demonstrated in low band gap non-lead containing
IV–VI semiconductor nanocrystals (such as SnTe), this would also represent a major
advancement with regards to these materials and their potential applications in PV
devices.

!
!
38
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!
49
Chapter 2. Synthesis and Characterization of Ternary Sn
x
Ge
1-x
Se
Nanocrystals*
*Published in Chem. Mater. 2012, 24, 3514–3516.

2.1. Introduction
The ability to tune IV–VI semiconductor nanocrystals through alloying can have a direct
effect on their optoelectronic properties, which can in turn affect device performance.
1-2

For example, Alivisatos et al. described the use of ternary PbS
x
Se
1-x
nanocrystals in
Schottky junction photovoltaic (PV) devices that demonstrated their compositionally
dependent device performance.
1
Schottky devices made with these PbS
x
Se
1-x

nanocrystals gave higher power conversion efficiencies (3.3%) than either of the PbS
(∼1.8%) or PbSe (∼1.4%) end-member lead chalcogenides. Although the potential utility
of alloyed ternary semiconductor nanocrystals is clear, their synthesis is often difficult.
3

Inherent differences in precursor reaction kinetics often lead to inhomogeneous internal
structures (i.e., gradient or core/shell structures).
4,5
A number of strategies have been
developed to balance the reactivity of the precursors and obtain homogeneous alloys in
the following IV–VI systems: PbSe
x
Te
1-x
, PbS
x
Se
1-x
, Pb
x
Sn
1-x
Te, PbS
x
Te
1-x
, Pb
x
Sn
1-x
S
Sn
x
Ge
1−x
S, Sn
x
Ge
1−x
Se GeS
x
Se
1−x
, and SnS
x
Se
1-x
,
4-13

The layered IV–VI semiconductor nanocrystals (i.e., SnS, SnSe, GeS, GeSe) have
recently gained popularity as potential alternatives to the lead chalcogenides, due in part
to their relatively higher stability and environmental sustainability.
14
In the bulk, SnSe is
!
!
50
a native p-type semiconductor with a high absorption coefficient (∼1 × 10
4
cm
–1
) and
narrow indirect and direct band gaps (E
g
= 0.9 and 1.3 eV, respectively) that are close to
the optimal band gap values for single junction solar cells.
15-17
Similar to the lead
chalcogenides, SnSe has a relatively large Bohr exciton radius (>23 nm) and exhibits
strong size-dependent optical properties enforced by quantum confinement effects.
18-19

However, with the exception of the work by Hanrath and co-workers, it has thus far been
difficult to tune the band gap of resulting SnSe nanocrystals through nanocrystal size
control within the quantum confined regime.
20-22

Another potential way of tuning the optoelectronic properties of the SnSe nanocrystals is
to alloy SnSe and GeSe to make ternary Sn
x
Ge
1-x
Se nanocrystals. Alloyed semiconductor
nanocrystals can display compositionally tunable properties, distinct from both their bulk
alloys and binary nanocrystalline end-members. Bulk GeSe exhibits similar properties to
SnSe. Both materials possess a thermodynamically preferred orthorhombic crystal
structure described as a highly distorted rock salt structure consisting of strongly bound
double layers.
23
In addition to being a native p-type semiconductor with a high
absorption coefficient, GeSe has a slightly larger band gap (E
g
= 1.1 – 1.2 eV) than SnSe
with reduced lattice parameters.
24
Previous to this work the SnxGe
1-
x Se alloy had only
been synthesized in the bulk form with complete solid solubility (0 ≤ x ≤ 1),
25-26
however,
since this publication additional reports on the synthesis of various tin and germanium
ternary chalcogenides have been reported.
13
In this chapter, we present the first synthesis
of ternary Sn
x
Ge
1-x
Se nanocrystals, and demonstrate tunable compositions throughout the
!
!
51
entire alloy range (0 ≤ x ≤ 1). Compositional tuning of the lattice parameters, band gaps,
and morphologies was demonstrated and the alloy formation mechanism was
investigated.
2.3. Results and Discussion
2.3.1. Synthesis of Sn
x
Ge
1-x
Se Nanocrystals
Both the ternary Sn
x
Ge
1-x
Se nanocrystals and the SnSe/GeSe end-members were
synthesized using a di-tert-butyl diselenide precursor, as previously reported by our
group for the synthesis of a variety of metal selenide nanocrystals.
27-29
In short, all
ternary Sn
x
Ge
1-x
Se nanocrystals were prepared by injecting di-tert-butyl diselenide (0.10
mmol) and hexamethyldisilazane (HMDS, 10 mmol) into a dodecylamine solution (13
mmol) containing varying ratios of GeI
4
and SnI
4
(0.40 mmol total) at 95 °C. The
reaction was then heated to 225 °C and held at that temperature for 4.75 h. It is
interesting to note that without the addition of HMDS, the reaction produces phase pure
SnSe nanocrystals with no evidence of germanium incorporation (Figure 2.1.). Therefore,
the addition of HMDS was found to be essential for alloy formation as it may aid in the
reduction of Ge(IV) to Ge(II), which is consistent with the more negative reduction
potential of Ge(IV) relative to Sn(IV). The SnSe and GeSe end-members were
synthesized using identical procedures to that described above with the exclusion of
either GeI
4
or SnI
4
, respectively. Additionally, obtaining crystalline GeSe in the
orthorhombic Pnma phase required the omission of HMDS and higher injection
!
!
52
temperatures (180 °C). When GeSe was synthesized in the presence of HMDS, an
unknown phase of germanium selenide was obtained with a 50:50 Ge/Se stoichiometry
(Figure 2.2.).

Figure 2.1. XRD pattern of nanocrystal product from a reaction containing a precursor Sn/Ge
ratio of 3:2, performed without the addition of HMDS. The pattern is equivalent to that obtained
when synthesizing SnSe using the current procedure and can be indexed to SnSe in the Pnma
space group.

Figure 2.2. XRD pattern of the unknown GeSe product synthesized with HMDS. EDS confirmed
a 1:1 Ge:Se stoichiometry.

!
!
53
2.3.2. Characterization of Ternary Sn
x
Ge
1-x
Se Nanocrystals
The elemental composition of the Sn
x
Ge
1-x
Se nanocrystals was determined using energy
dispersive X-ray spectroscopy (EDS). The composition of the ternary Sn
x
Ge
1-x
Se
nanocrystals was modulated by varying the stoichiometry of the SnI
4
and GeI
4
precursors
used during synthesis (Table 2.1., Figure 2.3.). A slight deviation between the Sn/Ge ratio
introduced during synthesis and the Sn/Ge ratio incorporated into the isolated Sn
x
Ge
1-x
Se
nanocrystals was observed. In compositions where x < 0.9, it was observed that greater
than nominal germanium concentrations were incorporated into the Sn
x
Ge
1-x
Se
nanocrystals at the expense of tin incorporation. Furthermore, slight stoichiometric
deviations from the expected 50:50 metal(loid)/chalcogen ratio were also observed in all
samples. All of the nanocrystals were selenium-poor (with deviations between 2.8 and
10%) given the selenium-limited (2:1 metal(loid)/selenium mole ratio) reaction
conditions required for alloy formation.
Table 2.1. Estimated compositions determined using EDS.
Sample Denotation Determined Sn:Ge:Se
ratio
SnSe 55:0:45
Sn
0.9
Ge
0.1
Se 47:5:48
Sn
0.7
Ge
0.3
Se 37:14:49
Sn
0.6
Ge
0.4
Se 30:22:48
Sn
0.4
Ge
0.6
Se 22:33:45
Sn
0.2
Ge
0.8
Se 11:44:45
Sn
0.1
Ge
0.9
Se 7:48:45
!
!
54

Figure 2.3. Elemental composition comparing the actual germanium incorporation in the isolated
nanocrystals (measured by EDS) versus expected from precursor addition.

Powder X-ray diffraction (XRD) revealed that all of the Sn
x
Ge
1-x
Se nanocrystals were
prepared phase pure (Figure 2.4.). Nanocrystals of both SnSe and GeSe were indexed to
the orthorhombic phase of the GeS structure, with a Pnma space group (PDF#01–075–
6133 and PDF#01–075–1802 for SnSe and GeSe, respectively). The calculated lattice
constants extracted from Rietveld analysis for SnSe (a = 11.5040(9) Å, b = 4.1563(4) Å,
c = 4.4301(5) Å) and GeSe (a = 10.841(2) Å, b = 3.8335(8) Å, c = 4.3877(12) Å) closely
match literature values. The Sn
x
Ge
1-x
Se nanocrystals display the same overall diffraction
patterns as the SnSe and GeSe end-members, indicating that they are isostructural with
the orthorhombic Pnma crystal structure. As expected with the smaller cationic radius of
Ge
2+
(cationic radii of Ge
2+
and Sn
2+
are 0.73 Å and 0.93 Å, respectively),
30
a gradual
shift to higher 2θ values was observed as the germanium content in the Sn
x
Ge
1-x
Se
!
!
55
nanocrystals increased. The lattice constants and unit cell volumes of the Sn
x
Ge
1-x
Se
nanocrystals were determined from the XRD patterns using Rietveld analysis (Figures
2.5. and 2.6., Table 2.2.) and compared against the experimentally determined elemental
compositions (Figure 1.1.b). The a and b lattice constants clearly demonstrate a linear
dependence on the composition as they decrease monotonically with increasing
germanium content. This linear relationship is consistent with Vegard’s law and
establishes the compositional homogeneity of the nanocrystals in these crystallographic
directions. Noteworthy, however, is the fact that the c lattice constant varies slightly from
linearity and tends to bow with compositional changes even though the overall unit cell
volume appears to vary linearly with composition. The bowing phenomenon of the c
lattice constant is in accord with previously published literature for bulk material.
25

Figure 2.4. Structural characterization of Sn
x
Ge
1–x
Se nanocrystals. (a) XRD patterns for select
Sn
x
Ge
1–x
Se compositions showing the orthorhombic Pmna crystal structure. (b) Dependence of
the lattice constants and unit cell volume on germanium content.
!
!
56

Figure 2.5. Rietveld analysis of XRD patterns of Ge
x
Sn
1−x
Se samples. Experimental (×) and
calculated (") patterns are shown for each sample along with the difference curve (") and
tickmarks (#) corresponding to the phase refined.

!
!
57

Figure 2.6. Rietveld analysis of XRD patterns of Ge
x
Sn
1−x
Se samples. Experimental (×) and
calculated (") patterns are shown for each sample along with the difference curve (") and
tickmarks (#) corresponding to the phase refined.

!
!
58!
Table 2.2. Rietveld Analysis of X-ray Diffraction Data of Ge
x
Sn
1−x
Se
x (mol. %)
a
0 10 30 40 60 80 90 100
a (Å) 11.5040(9) 11.4470(9) 11.3441(14) 11.2474(13) 11.1705(13) 11.0101(9) 10.9371(8) 10.841(2)
b (Å) 4.1563(4) 4.1244(4) 4.0667(6) 4.0158(6) 3.9813(5) 3.9001(5) 3.8727(4) 3.8335(8)
c (Å) 4.4301(5) 4.4365(4) 4.4374(7) 4.4348(8) 4.4251(6) 4.4023(6) 4.3967(6) 4.3877(12)
V (Å
3
) 211.82(4) 209.46(4) 204.73(6) 200.31(6) 196.80(5) 189.08(4) 186.23(4) 182.35(8)
X
Sn,Ge
0.1190(2) 0.1185(2) 0.1180(2) 0.1190(3) 0.1195(3) 0.1209(2) 0.1215(2) 0.1224(7)
Z
Sn,Ge
0.8994(5) 0.8979(4) 0.8927(6) 0.8894(6) 0.8882(5) 0.8914(5) 0.8917(5) 0.8847(13)
X
Se
0.8557(3) 0.8543(3) 0.8549(4) 0.8550(4) 0.8543(3) 0.8540(2) 0.8540(2) 0.8514(6)
Z
Se
0.5149(7) 0.5152(5) 0.5145(7) 0.5103(7) 0.5064(5) 0.5031(6) 0.5000(6) 0.4908(12)
U
Sn,Ge
(Å
2
)
b
2.37(12) 2.15(10) 2.12(15) 5.60(16) 3.27(13) 6.13(14) 6.20(14) 2.3(3)
U
Se
(Å
2
)
b
2.52(17) 2.38(14) 2.44(19) 5.34(19) 3.17(15) 4.98(14) 5.19(15) 0.4(3)
R
wp
8.7 6.2 9.2 8.0 6.8 4.9 6.5 13.0
a
Ge fraction extracted from energy dispersive X-ray analysis.
b
Given as 100×U.
!
!
59
Diffuse reflectance UV–vis–NIR spectroscopy was used to investigate the optical
properties of the Sn
x
Ge
1-x
Se nanocrystals. It was observed that as the composition of the
Sn
x
Ge
1-x
Se nanocrystals becomes more germanium rich, the onset of absorption
systematically blue-shifted from ca. 1400 nm (Figure 2.7.). Applying Kubelka–Munk
functions to the reflectance spectra allowed indirect band gap transitions to be estimated
([F(R)hν]
0.5
), which were found to range from E
g
= 0.87–1.13 eV as the composition of
the Sn
x
Ge
1-x
Se nanocrystals was tuned from x = 1.0 – 0.2, respectively (Figures 2.8. and
2.9.). These values generally fall within the range of bulk indirect band gaps previously
reported for SnSe and GeSe nanocrystals.
21,31

Figure 2.7. Diffuse reflectance spectra for Sn
x
Ge
1–x
Se nanocrystals. The spectra have been offset
for clarity.

!
60

Figure 2.8. Tauc plot of the linear regions from [F(R)hν]
1/2
as a function of energy for the
Sn
x
Ge
1-x
Se alloys showing the fits used to obtain the indirect band gap energies.

Figure 2.9. Calculated indirect band gaps ([F(R)hν]
0.5
) as a function of alloy composition.

!
61
Homogeneous nanocrystal alloys can result when the nucleation and growth rates of the
two constituent materials are similar. Here, it is apparent that SnI
4
is kinetically more
reactive than GeI
4
toward the active selenium source, as evidenced by a shorter
nucleation time for the SnSe nanocrystal synthesis (i.e., <20 min as compared to >1 h for
the GeSe nanocrystal synthesis). This obvious reactivity difference prompted an
investigation into the mechanism of alloy formation. Timed-aliquot XRD studies were
performed to monitor the change in nanocrystal composition during growth. For example,
when SnI
4
and GeI
4
were added to the reaction in a 4:1 mol ratio, the initial XRD pattern
of the aliquot taken ca. 20 min after nucleation appeared close to that of SnSe (a = 11.44
Å, b = 4.16 Å, c = 4.35 Å). Subsequent XRD patterns of aliquots taken at longer time
intervals display a gradual shift to higher 2θ values as the reaction proceeds to
completion. After 4.75 h, the nanocrystals exhibit lattice constants of a = 11.28 Å, b =
4.04 Å, and c = 4.43 Å, as expected from nanocrystals with an approximate Sn
0.6
Ge
0.4
Se
composition (vide supra; Figures 2.10. and 2.11.). This implies that nucleation of the
crystalline material begins as SnSe (or a tin-rich selenide) and gradually incorporates
germanium over the growth period, suggesting the potential involvement of a cation
exchange mechanism. This also points to the difference in the nucleation and growth
kinetics, where the nucleation kinetics under the current conditions favor the formation of
SnSe, whereas germanium incorporation or exchange occurs during growth until the final
composition is achieved.
!
62

Figure 2.10. Timed aliquot XRD patterns taken when using a SnI
4
/GeI
4
ratio of 4:1. A gradual
shift to higher 2θ values is observed as the reaction proceeds to completion (ca.150 min). After
4.75 h, the XRD pattern matches that expected from nanocrystals with an approximate
Sn
0.6
Ge
0.4
Se composition.

Figure 2.11. Lattice parameter a as a function of time taken from timed aliquot XRD patterns
when using a SnI
4
/GeI
4
ratio of 4:1.

2.3.3. Formation Mechanism of Sn
x
Ge
1-x
Se Nanocrystals
To investigate the possibility of a cation exchange mechanism, we first synthesized SnSe
nanocrystals and then allowed them to react with the germanium source via injection of
!
63
GeI
4
in dodecylamine after 1 h. Powder XRD and EDS were used to monitor the
structural and compositional changes before and after addition of the GeI
4
to the SnSe
nanocrystals. The XRD pattern taken before introduction of GeI
4
matches that of SnSe,
while the XRD pattern taken 3.75 h after introduction of GeI
4
suggests germanium
incorporation with lattice constants of a = 11.14 Å, b = 3.96 Å, and c = 4.43 Å, in
addition to some unreacted SnSe (Figure 2.12.). EDS confirms the compositions before
and after introduction of GeI
4
as nearly 50:50 Sn/Se and 25:25:50 Sn/Ge/Se, respectively.
These results suggest that germanium may indeed alloy into SnSe through a possible
cation exchange mechanism. This is similar to a recent report by Schaak and co-workers
that described the transformation of SnSe to SnTe through an anion exchange
mechanism.
32

Figure 2.12. XRD patterns of the nanocrystal products taken when first synthesizing SnSe (blue
pattern) and after injection of GeI
4
to the SnSe nanocrystal product (total Sn/Ge ratio of 3:2,
green pattern). The XRD pattern of the resulting nanocrystals matches closely to that obtained in
the typical alloy synthesis when using a SnI
4
/GeI
4
ratio of 3:2, with a small fraction of
unconverted SnSe.
!
64
2.3.4. Morphological Dependence on Composition
Interesting morphological changes also resulted as the Sn
x
Ge
1-x
Se nanocrystal
composition was varied. Transmission electron microscopy (TEM) was used to monitor
the morphological differences between the sheetlike SnSe nanocrystals and the rodlike
morphologies exhibited by the germanium-rich Sn
0.1
Ge
0.9
Se nanocrystals (Figure 2.13.).
The SnSe nanosheets adopted rectangular morphologies with lateral dimensions varying
between 0.7–1.0 µm × 0.5–1.0 µm in size. The Sn
0.1
Ge
0.9
Se nanorods vary in size with a
range of dimensions from 70–100 nm in width to 0.5–1.5 µm in length. All intermediate
compositions inherited morphological features from both compositional extremes;
forming intermediate sheetlike aggregates comprised of bundled nanorods. This
phenomenon has been previously reported in Pb
x
Sn
1-x
S, where the PbS and SnS end-
members exhibit cube and spherical morphologies, respectively, and the intermediate
Pb
0.85
Sn
0.15
S nanocrystals inherit aspects of both end-members with truncated cube
morphologies.
8

Figure 2.13. TEM images of SnxGe1-xSe nanocrystals showing the compositionally dependent
morphology change for: (a) SnSe, (b) Sn
0.9
Ge
0.1
Se, (c,d) Sn
0.7
Ge
0.3
Se, (e) Sn
0.6
Ge
0.4
Se, (f)
Sn
0.4
Ge
0.6
Se, (g) Sn
0.2
Ge
0.8
Se, and (h) Sn
0.1
Ge
0.9
Se.
!
65
2.4. Experimental
2.4.1. Materials and Methods
Tin(IV) iodide (SnI
4
, Alfa Aesar, 95%), germanium(IV) iodide (GeI
4
, Strem, 99.999%)
and hexamethyldisilazane (HMDS, Aldrich, ≥99.0%) were all purchased and used
without further purification. Dodecylamine (Alfa Aesar, ≥98%) was distilled from CaO
prior to use. Nanocrystal syntheses were performed under nitrogen, in the absence of
water and oxygen, using standard Schlenk techniques.
Energy dispersive X-ray spectroscopy (EDX) spectra was collected on a JEOL JSM-6610
scanning electron microscope operating at 20 kV and equipped with an EDAX Apollo
silicon drift detector (SDD). Multiple regions of a sample deposited on an aluminum stub
were analyzed and averaged. Conventional powder X-ray diffraction (XRD) patterns
were collected in the 10–80° 2θ range using a Rigaku Ultima IV diffractometer operating
at 44 mA and 40 kV. Cu Kα radiation (λ = 1.5406 Å) was employed. Diffraction patterns
were recorded at 25 ̊C. Rietveld structural refinements were carried out using the General
Structure Analysis System (GSAS) software.
33
Experimental data and atomic X-ray
scattering factors were corrected for sample absorption and anomalous scattering,
respectively. The Sn:Ge molar ratio was fixed as determined by energy dispersive X-ray
analysis. The following parameters were refined: (1) scale factor, (2) background, which
was modeled using a shifted Chebyschev polynomial function, (3) peak shape, which was
modeled using a modified Thomson−Cox−Hasting pseudovoight function including an
!
66
asymmetry parameter, (4) lattice constants, (5) fractional atomic coordinates, (6) an
isotropic thermal parameter for each chemical species (i.e., USn,Ge, and USe), and (7)
preferred orientation, which was modeled using the March-Dollase approach. The usual
Rwp indicator was employed to assess the quality of the refined structural models.
Transmission electron microscopy (TEM) was performed on a JEOL JEM- 2100
microscope at an operating voltage of 200 kV, equipped with a Gatan Orius CCD camera.
Samples were prepared from dilute purified dispersions in toluene and deposited onto 300
mesh Formvar-coated copper grids (Ted Pella, Inc.). Diffuse reflectance measurements
were collected on a Perkin-Elmer Lambda 950 equipped with a 150 mm integrating
sphere. Samples were prepared by drop casting a solution of the nanocrystals in toluene
on a glass substrate to form a thick film. The sample was then covered by a quartz
window and the spectra were measured in reflectivity mode.
2.5. Conclusions
In summary, a facile synthetic method for compositional control in SnxGe
1–
xSe
nanocrystals has been described for the first time. Homogeneous internal structures were
obtained, despite the inherent relative differences in the cation precursor reaction rates.
Investigation into the formation mechanism revealed the initial formation of SnSe or tin-
rich selenide followed by a partial cation exchange with Ge
2+
. The band gaps, lattice
parameters, and morphologies of the alloys were tuned via nanocrystal composition. This
alloying approach provides an additional means of band gap tuning that is otherwise
!
67
difficult to achieve in the layered IV–VI semiconductors; making these materials
potential candidates as photovoltaic materials.
2.6. References
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Employing Ternary PbS
x
Se
1-x
Nanocrystals. Nano Lett. 2009, 9, 1699-1703.
(2) Yu, K.; Ouyang, J.; Zhang, Y.; Tung, H. T.; Lin, S.; Nagelkerke, R. A. L.;
Kingston, D.; Wu, X.; Leek, D. M.; Wilkinson, D.; Li, C.; Chen, I. G.; Tao, Y.
Low-Temperature Noninjection Approach to Homogeneously-Alloyed PbSe
x
S
1-x

Colloidal Nanocrystals for Photovoltaic Applications. ACS Appl. Mater.
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Compositionally Tunable Ternary PbSe
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as Well as Binary PbSe and PbS
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Gardonio, S.; O’Brien, P. Controlled Synthesis of Tuned Bandgap
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x
Se
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Synthesis of Alloy PbS
x
Se
1−x
Nanocrystals. J. Mater. Chem. 2012, 22, 5984-5989.
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1−x
Sn
x
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Ternary Alloy Nanocrystals of Tin and Germanium Chalcogenides. RSC Adv.
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Monochalcogenide IV–VI Semiconductor Nanocrystals for Use in Solar Cells.
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Films. Appl. Surf. Sci. 2011, 257, 2944-2949.
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of Quantum Sized Semiconducting Tin(II) Selenide in Thin Film Form. Thin
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Synthesis of Unprotected Nanoparticles of SnTe and SnSe and a New Route to
(PhSe)
4
Sn. Inorg. Chem. 2002, 41, 6001-6005.
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Phase Synthesis of SnSe Nanocrystals for Use in Solar Cells. J. Am. Chem. Soc.
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119 Mössbauer and Electrical Conductivity Study of the System Sn
x
Ge
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Se (0 ≤ x
≤ 1) Solid State Commun. 1978, 27, 1185-1187.
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(27) Norako, M. E.; Brutchey, R. L. Synthesis of Metastable Wurtzite CuInSe
2

Nanocrystals. Chem. Mater. 2010, 22, 1613-1615.
(28) Webber, D. H.; Brutchey, R. L. Photochemical Synthesis of Bismuth Selenide
Nanocrystals in an Aqueous Micellar Solution. Inorg. Chem. 2011, 50, 723-725.
(29) Norako, M. E.; Greaney, M. J.; Brutchey, R. L. Synthesis and Characterization of
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Colloidal Nanosheets of GeS and GeSe. J. Am. Chem. Soc. 2010, 132, 15170-
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R. E. Engineering Porosity into Single-Crystal Colloidal Nanosheets Using
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Alamos National Laboratory, 2000.

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71
Chapter 3. Chalcogenol Ligand Toolbox for CdSe Nanocrystals and Their
Influence on Exciton Relaxation Pathways*
*Published in ACS Nano 2014, 8, 2512–2521.

3.1. Abstract
We have employed a simple modular approach to install small chalcogenol ligands on the
surface of CdSe nanocrystals. This versatile modification strategy provides access to
thiol, selenol, and tellurol ligand sets via the in situ reduction of R
2
E
2
(R =
t
Bu, Bn, Ph; E
= S, Se, Te) by diphenylphosphine (Ph
2
PH). The ligand exchange chemistry was
analyzed by solution NMR spectroscopy, which reveals that reduction of the R
2
E
2

precursors by Ph
2
PH directly yields active chalcogenol ligands that subsequently bind to
the surface of the CdSe nanocrystals. Thermogravimetric analysis, FT-IR spectroscopy,
and energy dispersive X-ray spectroscopy provide further evidence for chalcogenol
addition to the CdSe surface with a concomitant reduction in overall organic content from
the displacement of native ligands. Time-resolved and low temperature
photoluminescence measurements showed that all of the phenylchalcogenol ligands
rapidly quench the photoluminescence by hole localization onto the ligand. Selenol and
tellurol ligands exhibit a larger driving force for hole transfer than thiol ligands and
therefore quench the photoluminescence more efficiently. The hole transfer process could
lead to engineering long-lived, partially separated excited states.

!
72
3.2. Introduction
The optoelectronic properties of colloidal semiconductor nanocrystals are largely
dependent on their surface chemistry as a result of their large surface-to-volume ratios.
1-4

While the native ligands that are installed during nanocrystal synthesis impart solution
dispersibility and may passivate surface trap states, these ligands are typically long-chain
aliphatic compounds that are also electrically insulating (e.g., C18 fatty acids like stearic
acid), and they therefore have an extremely detrimental effect on conductivity and charge
mobility in nanocrystal-based thin film devices.
5,6
As a result, there is currently a
tremendous amount of interest in controlling the surface chemistry of colloidal
semiconductor nanocrystals via ligand exchange approaches.
5,7,8
In such an approach,
the native ligands are exchanged with less sterically demanding ligands that ideally still
impart solution dispersibility and effectively passivate the nanocrystal surface, but which
also allow for more efficient interparticle charge transfer in nanocrystal thin film devices.
The CdSe nanocrystal surface is generally cadmium-rich,
9,10
which necessitates charge
balance with an anionic X-type ligand.
11,12
Much success has been had with ligand
exchanges that employ both strong and weakly coordinating inorganic X-type anions,
such as chloride,
13
thiocyanate,
14
sulfide,
15-17
chalcogenometallates,
18
and
tetrafluoroborate;
19,20
however, less success has been achieved with small organic
ligands. Traditionally, organic ligand exchanges have been done with pyridine and small
primary amines,
21-24
but such neutral L-type ligands cannot exchange the strongly bound
X-type ligands on nonstoichiometric nanocrystal surfaces unless there is some degree of
!
73
surface reconstruction (e.g., through loss of CdX
2
).
11
We previously reported that tert-
butylthiol could be employed as a small organic ligand to quantitatively displace native
long-chain carboxylate and phosphonate ligands on CdSe nanocrystals and yield
colloidally stable dispersions.
25
After ligand exchange with tert-butylthiol, a significant
decrease in organic content was observed by thermogravimetric analysis (TGA) and FT-
IR spectroscopy. Importantly, ligand exchange with tert-butylthiol leads to significantly
improved photocurrent, charge mobility, and power conversion efficiencies in
nanocrystal-based devices relative to those using nonligand-exchanged or pyridine-
exchanged CdSe nanocrystals.
25-27
In addition to affecting quantitative ligand exchange
of both X-type and L-type ligands, it is known that soft thiol ligands bind very strongly to
the soft cadmium atoms on the surface of CdSe.
28,29
From this, it follows that the
analogous selenols and tellurols would also be expected to have a strong binding affinity
for CdSe nanocrystal surfaces as a result of their chemically soft nature. To the best of
our knowledge, however, there have been no studies on the binding of organotellurol
ligands to CdSe nanocrystals, and only one descriptive example of phenylselenolate
binding to CdSe clusters using PhSeSiMe
3
in 1988.
30
Thus, there is interest in
developing simple, modular approaches for the installation of selenol and tellurol ligands
onto CdSe nanocrystals such that the fundamental ligand exchange chemistry and
photophysics may be studied on these systems and compared to that of traditional thiol
ligands.
!
74
The photophysical study of nanocrystals offers an experimental window into their
electronic structure and into the fate of photogenerated charge carriers. Energetic
positions of the conduction and valence band edges are determined by the material and
size of the nanocrystal, due to quantum confinement effects, resulting in tunable
absorption and emission properties. The positions of the band edge levels relative to that
of the surrounding molecules or lack thereof (e.g., surface states created by dangling Cd
or Se bonds) also affect the oxidative and reductive abilities of the nanocrystals.
Positively or negatively doping nanocrystals via charge transfer to the ligands modifies
the emission characteristics of the nanocrystals by competing with other radiative and
nonradiative pathways, and conversely, the nanocrystal emission offers an opportunity to
study charge transfer processes and dynamics. The influence of thiol ligands on the
photophysical properties of CdSe nanocrystals has been extensively studied; transient
absorption and photoluminescence (PL) studies have shown that thiol and thiolate ligands
act as photogenerated hole traps and passivate electron traps.
31,32
On the other hand, the
influence of selenol and tellurol ligands is much less explored to date. Recently, PbSe
nanocrystals with a native oleate ligand shell were subjected to a ligand exchange with
long-chain octyldecylselenol, which was shown to bind to the nanocrystal surface as
selenolate through a proton transfer mechanism with release of free oleic acid.
33

Investigation into the photophysical properties resulting from octyldecylselenol treatment
on PbSe revealed the introduction of a new nonradiative decay pathway, identified as a
hole-trap; similar effects have been observed after selenide (Se
2–
) treatment of CdSe.
3
!
75
In this chapter, we report a simple and modular ligand exchange procedure to replace
native stearate ligands on CdSe nanocrystals using a full series of seven small, air-stable
R
2
E
2
dichalcogenide precursors (R = tBu, E = Se or Te; R = Bn, E = S or Se; R = Ph, E =
S, Se, or Te). This ligand exchange is accomplished by the reduction of the R
2
E
2

dichalcogenide with diphenylphosphine (Ph
2
PH) to generate the small chalcogenol ligand
in situ in a CdSe nanocrystal suspension. Upon workup, this approach yields colloidally
stable, thiol-, selenol-, or tellurol-exchanged nanocrystals whose low temperature and
time-resolved photophysics were studied and compared. We show that four processes
compete for the relaxation of photoexcitations: (i) band-edge emission, (ii) phonon-
assisted nonradiative relaxation processes, (iii) surface state trapping, and finally, (iv)
hole trapping by phenylchalcogenol ligands. While the importance of these processes
depends on the temperature and ligand type, our findings show hole trapping to be more
efficient for the phenylselenol and phenyltellurol ligands than for phenylthiol. For these
ligands, the hole trapping mechanism could outcompete relaxation to the ground state via
surface state trapping. This finding could help in the engineering of long-lived, partially
separated excited states.
3.3. Results and Discussion
3.3.1. Colloidal Ligand Exchange
The CdSe nanocrystals were synthesized via the hot injection of tri-n-octylphosphine
selenide (TOPSe) into a solution of Cd(stearate)
2
in tri-n-octylphosphine oxide (TOPO,
!
76
98%) and technical grade stearic acid.
34
The resulting nanocrystals were purified three
times using toluene as a dispersant and ethanol as a flocculant, with the final product
being dispersed in toluene. The UV–vis spectrum of the resulting nanocrystals dispersed
in toluene displayed a clear first exciton peak at 586 nm (2.12 eV), which is indicative of
a CdSe nanocrystal diameter of 4.0 nm.
35
Transmission electron microscopy (TEM)
revealed the CdSe nanocrystals to possess a spherical nanoparticle morphology with
diameters of 3.9 ± 0.5 nm, which is in close agreement with empirical sizing by UV–vis
spectroscopy.
The nanocrystals with the native ligand shell, CdSe(NL), can be ligand exchanged with
seven R
2
E
2
dichalcogenide precursors (R =
t
Bu, E = Se or Te; R = Bn, E = S or Se; R =
Ph, E = S, Se, or Te) using a facile, one-step, room temperature reduction with Ph
2
PH to
generate the corresponding chalcogenol. Addition of R
2
E
2
(0.30 mmol) and Ph
2
PH (0.92
mmol) to a stirred suspension of CdSe(NL) nanocrystals in toluene (1.0 mL at a
gravimetrically determined concentration of 17 mg mL
–1
corrected to pure CdSe content
based on the 500 °C TGA mass) resulted in ligand exchange, as indicated by nanocrystal
precipitation from the nonpolar solvent system. Ligand exchange reactions were
performed under nitrogen using standard Schlenk techniques; however, the nanocrystals
were handled in air post ligand exchange. Large differences in the reaction rates were
observed for the ditellurides, diselenides, and disulfides. Addition of the ditellurides (i.e.,
tBu
2
Te
2
and Ph
2
Te
2
) in the presence of Ph
2
PH to a CdSe(NL) suspension resulted in
immediate nanocrystal precipitation, the diselenides (i.e., tBu
2
Se
2
, Bn
2
Se
2
, and Ph
2
Se
2
)
!
77
required a several hours (ca. 4 h) before nanocrystal precipitation, and the disulfides (i.e.,
Bn
2
S
2
and Ph
2
S
2
) gradually precipitated the nanocrystals over a 24 h period. Removal of
excess precursors, reaction byproducts, and displaced native ligands required a total of
four washes using tetramethylurea (TMU) as the dispersant and pentane as a flocculant.
The dispersibility of the CdSe nanocrystals was found to drastically change after ligand
exchange. The CdSe(NL) nanocrystals are soluble in low polarity solvents, such as
toluene and hexanes, as a result of the long-chain aliphatic native ligands. Replacement
of the native ligands for smaller tert-butyl-, benzyl-, or phenylchalcogenols causes the
nanocrystals to lose colloidal stability in nonpolar solvents; however, after ligand
exchange the CdSe nanocrystals are readily dispersible in moderately polar solvents, such
as TMU. Nanocrystal suspensions in TMU maintained high colloidal stability (>3
months), were optically transparent and were easily passed through a 100 nm filter.
Additionally, it should be noted for optical studies that the line width of the first exciton
peak remains unchanged after ligand exchange and washing indicating that the
nanocrystal size population is similarly unchanged. While TMU proved to be the best
solvent for these systems, they were also readily dispersible in other donor solvents, such
as dimethyl sulfoxide (DMSO), N,N-dimethylacetamide (DMAc), and pyridine. The
ligand exchanged CdSe nanocrystals were also dispersible in less polar solvents, such as
1,2-dichlorobenzene.

!
78
3.3.2 NMR Characterization
Multinuclear
1
H,
31
P, and
77
Se NMR spectroscopy was used to gain insight into the nature
of the native ligands and their ligand exchange with the chalcogenols generated by in situ
reduction. The conditions used to synthesize the CdSe nanocrystals in this study are
known to produce native ligand shells primarily comprising stearate.
25
To verify this,
NMR spectra were recorded for both the surface-bound native ligands and free native
ligands from acid digested nanocrystals. The
1
H NMR spectrum of the purified
CdSe(NL) nanocrystals in a CDCl
3
suspension displayed only the characteristic
resonances of bound stearate ligands at 0.94, 1.28, and 1.60 ppm (Figure 3.1.). These
NMR resonances display the typical broadening and downfield chemical shifts of surface
bound species, with the α-methylene protons of the carbonyl group being essentially
unobservable as a result of their proximity to the binding headgroup.
36-38
The native
ligands were further analyzed after a typical digestion and extraction procedure using
aqua regia.
18
The
1
H and
31
P NMR spectra in benzene-d
6
of the liberated ligands suggest
that the CdSe(NL) surface is free of any detectable alkylphosphonic acid species and is
ligated solely by stearic acid (δ = 0.92, 1.35, 1.47, 2.04 ppm) (Figure 3.2.).

!
79

Figure 3.1.
1
H NMR spectrum of the purified CdSe(NL) particles in CDCl
3
showing only the
characteristic resonances of bound stearate ligands, which are significantly broadened and shifted
downfield (remaining peaks due to residual solvents toluene δ = 7.24, 7.16 and 2.36 ppm and
methanol δ = 3.49 ppm). No free ligands are observed.

Figure 3.2.
1
H (top) and
31
P (bottom) NMR spectra of the digestion products from CdSe(NL)
nanocrystals in benzene-d
6
.

The ligand exchange chemistry was monitored using multinuclear NMR spectroscopy;
the stepwise addition of the reagents to benzene-d
6
elucidated the reduction and ligand
exchange with Ph
2
Se
2
as a model system (Figure 3.3.). The reaction between Ph
2
Se
2
and
Ph
2
PH (25 °C, 24 h) produced two products shown in Figure 3.3.a,b, unambiguously
identified as Ph
2
PSePh (
77
Se δ = 308.8 ppm;
31
P δ = 29.9 ppm (d,
1
J
P–Se
= 229 Hz); m/z =
!
80
341.99) and PhSeH (
77
Se δ = 145.5 ppm;
1
H δ = 1.26 (s), 7.16 (m) ppm; m/z = 157.99) by
NMR spectroscopy and mass spectrometry.
39
Subsequent addition of CdSe(NL)
nanocrystals to the same reaction of Ph
2
Se
2
and Ph
2
PH in benzene-d
6
(25 °C, 24 h)
resulted in binding of the in situ generated PhSeH, as evidenced by the near complete
disappearance of the sharp selenol resonances (
77
Se δ = 145.5 ppm;
1
H δ = 1.26 ppm) and
the broadening of the associated aromatic
1
H NMR signal at δ = 7.17 (Figures 3.3.a,b, 3.4
– 3.6.). The appearance of a broad resonance (i.e.,
1
H δ = 2.7–3.0 ppm) that is present
upfield from the unbound selenol Se–H resonance also suggests binding of the selenol
ligand (Figures 3.3.a and 3.7.). Analogous NMR studies involving the addition of free
thiols to CdSe nanocrystals have reported similar resonances for the sulfhydryl proton
that appear broadened and shifted upfield from the free ligands; these signature
resonances are attributed to bound thiols.
11
In addition, the release of a small amount of
free stearic acid was observed by
1
H NMR spectroscopy (Figure 3.7.), presumably
resulting from a proton transfer exchange mechanism occurring between the selenol and
stearate on the CdSe nanocrystal surface.
33
Given the stronger acidic character of
phenylselenol compared to stearic acid (pK
a
= 10.15 for stearic acid in H
2
O; pK
a
= 5.9 for
phenylselenol in H
2
O), the equilibrium is expected to shift toward the formation of the
selenolate.
40,41
Therefore, the NMR evidence suggests that the ligand may be binding as
mixture of neutral L-type selenol and anionic X-type selenolate to the CdSe nanocrystal
surface, as was previously proposed for the tert-butylthiol ligand.
25

!
81

Figure 3.3. (a)
1
H and (b)
77
Se NMR spectra of the reaction between Ph
2
Se
2
and Ph
2
PH before
and after the addition of CdSe(NL) nanocrystals. Full
1
H NMR spectra from 0-8 ppm are
provided in Figure 3.4.

!
82

Figure 3.4.
1
H NMR spectra monitoring the reactions resulting from sequential addition of
Ph
2
Se
2
(0.75 mmol) (bottom) to 1 equivalent of Ph
2
PH (middle) to CdSe(NL) (top) in 0.6 mL
benzene-d
6
. The CdSe(NL) nanocrystals were minimally dispersed in a small volume of
benzene-d
6
. The new
1
H peak (~3 ppm) in the top spectrum is residual MeOH (CH
3
resonance)
used to flocculate the CdSe(NL) nanocrystals. Reactions were run overnight at 25 ºC before
spectra were taken.

!
83

Figure 3.5.
77
Se NMR spectra of the reaction described in Figure 3.4. Reactions resulting from
sequential addition of Ph
2
Se
2
(0.75 mmol) (bottom) to 1 equivalent of Ph
2
PH (middle) to
CdSe(NL) (top) in 0.6 mL benzene-d
6
. The CdSe(NL) nanocrystals were minimally dispersed in
a small volume of benzene-d
6
. Reactions were run overnight at 25 ºC before spectra were taken.

!
84

Figure 3.6.
31
P NMR spectra of the reaction described in Figure 3.4. Reactions resulting from
sequential addition of Ph
2
Se
2
(0.75 mmol) (bottom) to 1 equivalent of Ph
2
PH (middle) to
CdSe(NL) (top) in 0.6 mL benzene-d
6
. The CdSe(NL) nanocrystals were minimally dispersed in
a small volume of benzene-d
6
. Reactions were run overnight at 25 ºC before spectra were taken.

!
85

Figure 3.7. Magnified
1
H NMR spectrum of the reaction between Ph
2
Se
2
, Ph
2
PH and CdSe(NL)
nanocrystals displaying the low intensity resonances of the bound selenol and free stearic acid.

3.3.3. Efficacy of Ligand Exchange
Thermogravimetric analysis was used as one means of determining the degree of ligand
exchange on the CdSe nanocrystal surface. The TGA traces (ambient to 500 °C, 10 °C
min–1, under flowing nitrogen) of the ligand exchanged nanocrystals previously dried at
100 °C show a clear reduction in the total organic mass content when compared to the
CdSe(NL) nanocrystals (Figure 3.8.a and 3.9.). By approximating the mass at 500 °C to
be pure CdSe, it was found that dried CdSe(NL) nanocrystals contained 28% organic
content after three toluene/ethanol washes. Ligand exchange reduces the total organic
content by half (on average) and produces a new lower temperature (<300 °C) mass loss
event, which likely corresponds to loss of the new chalcogenol ligands.
25
To quantify the
extent of ligand exchange by TGA, it was assumed that mass loss events >300 °C
!
86
correspond to the tightly bound native stearate ligands (based on the TGA trace for the
CdSe(NL) nanocrystal sample and previous work), while all new mass loss occurring
<300 °C was attributed to the new chalcogenol ligands.
25
It should be noted that the
chalcogenol ligands may also contribute to mass loss >300 °C; however, the majority of
chalcogenol decomposition was assumed to be <300 °C for this analysis. Corroborating
analyses show this to be a reasonable approximation (vide infra). For the CdSe(NL)
nanocrystals, there is 0.39 g organics/g CdSe. The organic component is comprised of
stearate, as discussed previously, with a mass loss event between 300–500 °C. For
PhSH-exchanged CdSe nanocrystals, the 300–500 °C mass loss represents 0.08 g per
gram of CdSe, corresponding to an 80% reduction in the stearate mass content. In the
case of the PhSeH- and PhTeH-exchanged CdSe nanocrystals, the native ligand content is
reduced by 70 and 79%, respectively. The mass loss below 300 °C is attributed to the
new phenylchalcogenol ligands and represents 0.06 g per gram of CdSe for PhSH, 0.07 g
per gram of CdSe for PhSeH, and 0.08 g per gram of CdSe for PhTeH ligands. The
relative mass contributions of native stearate and of new ligands after ligand exchange
with the entire series of dichalcogenide precursors studied are given in Table 3.1.
!
87

Figure 3.8. Thermogravimetric and spectroscopic evidence for displacement of native ligands on
the CdSe nanocrystal surface upon introduction of phenylchalcogenols. (a) TGA traces, (b) FT-
IR spectra, (c) EDX spectra, and (d) UV-vis spectra of CdSe(NL) nanocrystals before (green
data) and after ligand exchange with PhSH (red data), PhSeH (blue data), and PhTeH (purple
data).

Figure 3.9. TGA traces of ligand exchange reactions with Bn
2
S
2
, Bn
2
Se
2
,
t
Bu
2
Se
2
, and
t
Bu
2
Te
2
precursors.

!
88
Table 3.1. Low temperature and high temperature mass loss percentages used to approximate
extent of ligand exchange.
Mass loss % NL PhSH PhSeH PhTeH BnSH BnSeH
t
BuSeH
t
BuTeH
< 300 ºC 0 5 6 7 12 11 9 8
> 300 ºC 28 7 10 7 2 6 3 2
Total % 28 12 16 14 14 17 12 10

FT-IR spectroscopic analysis of the CdSe nanocrystals before and after ligand exchange
with the phenylchalcogenols further validates the TGA data by showing large reductions
in the native ligand content and the appearance of new bands corresponding to aromatic
ν(C–H) stretches, as shown in Figure 3.8.b. Quantitative comparisons were made
possible through the use of an internal standard. All nanocrystal samples were dried at
100 °C before analysis, and all spectra were normalized in intensity to the 2089 cm
–1

ν(C≡N) stretching peak of a Fe
4
[Fe(CN)
6
]
3
internal standard. The normalized integrated
areas under the aliphatic ν(C–H) stretching region were used to quantify removal of
native ligands after ligand exchange since the aromatic stretching bands associated with
the phenylchalcogenol ligands at ∼3050 cm
–1
are well resolved from the ν(C–H) stretches
associated with the native ligands in the range of 2800–3000 cm
–1
. Ligand exchange
using the Ph
2
S
2
, Ph
2
Se
2
, and Ph
2
Te
2
dichalcogenide precursors resulted in a 79%, 73%,
and 84% reduction in native ligand content, respectively; these data are in good
agreement with TGA results. Reductions in the native ligand ν(C–H) stretching intensity
were also observed after ligand exchange with the Bn
2
E
2
and tBu
2
E
2
dichalcogenide
precursors, but quantification was not performed because the tert-butyl and benzyl groups
of the entering ligands possess aliphatic C–H bonds. In general, however, it was
!
89
qualitatively observed by FT-IR spectroscopy that these tert-butyl and benzylchalcogenol
ligands also successfully displaced a large percentage of the native ligands, which is
consistent with TGA results. Energy dispersive X-ray spectroscopy (EDX) further
corroborated the presence of sulfur and tellurium after exchange with the PhSH and
PhTeH ligands, respectively (Figure 3.8.c). In contrast, the CdSe(NL) nanocrystals do not
possess any sulfur or tellurium by EDX.

Figure 3.10. TEM images of CdSe(NL) and Ph
2
Te
2
exchanged particles.

TEM analysis also provides evidence for ligand exchange when utilizing the heavier, and
therefore more “visible” or electron-dense, tellurol ligands (Figure 3.10.). As previously
stated, the average diameter of the CdSe(NL) nanocrystals measured in ImageJ from
representative TEM micrographs was 3.9 ± 0.5 nm. After ligand exchange with the
PhTeH or tBuTeH ligands, it was evident that the mean CdSe nanocrystal diameter had
increased. Statistical analysis, using a nonparametric Wilcoxon rank sum test
(necessitated by the non-normal population distribution), of the nanocrystal diameters
both before and after ligand exchange confirmed the increase in diameter (n = 1001; p =
!
90
10
–142
). Size histograms comparing the two sets of nanocrystals (i.e., native ligand and
tellurol-exchanged) help to visualize the increase by showing a clear shift to larger
diameters for the tellurol-exchanged CdSe nanocrystals (Figure 3.11.). To further verify
the statistical significance of the size increase from ligand exchange, a box plot
representation of the data clearly shows that the notched regions do not overlap,
indicating with 95% confidence that the true median diameters are significantly different.
Quantification of the increase showed that after ligand exchange with PhTeH, the mean
diameter of the CdSe nanocrystals increased by ∼0.7 nm to 4.6 ± 0.5 nm. Likewise,
ligand exchange with tBuTeH resulted in an increase in mean diameter by ∼0.9 nm to 4.8
± 0.8 nm. The increase in nanocrystal size after addition of the tellurol ligands is roughly
consistent with the addition of a monolayer of tellurium to the nanocrystal surface,
assuming an ionic radius of 0.21 nm for Te
2–
(i.e., 4 × 0.21 nm =0.84 nm).

Figure 3.11. (Right) Histograms of nanocrystal diameters before (green) and after
(purple) ligand exchange with Ph
2
Te
2
. The diameters were measured in ImageJ from
TEM micrographs of the respective nanocrystals. The number of bins was defined by the
square root of the total number of diameters (n = 1001). (Left) Box plot of nanocrystal
!
91
diameters before (green) and after (purple) ligand exchanged with Ph
2
Te
2
. The diameters
are the same as measured in the histograms. The notched regions do not overlap,
indicating with 95% confidence that the true median diameters are significantly different.
The horizontal lines with arrows represent the median values of the measured diameters
and the top and bottom of each box represents the 75th and 25th percentiles, respectively.
The whiskers extend to include approximately 99.3% of the data and the outliers are
marked with crosses.

3.3.4. Photophysical Measurements
Due to their widespread use, the influence of thiols on the photophysical properties of
CdSe nanocrystals has been thoroughly investigated. The current study employs PhSH as
a benchmark while investigating the effect of the lesser-known selenol and tellurol
ligands on the photophysical properties of CdSe nanocrystals. In colloidal nanocrystals,
several processes are known to induce relaxation to the ground state. They have been
identified as radiative relaxation (band-edge emission), Auger nonradiative scattering,
42

Förster energy transfer to bigger nanocrystals, thermal escape,
43-45
trapping in surface
and/or defects states,
32,43,46-49
and ligand-induced charge transfer.
31,32,46,49,50
The large
number of identified processes, together with the intrinsic size and surface chemistry
distribution of samples, typically results in complex decay profiles. Here, by performing
time-resolved photoluminescence measurements at low excitation densities,
concentrations, and temperatures, we reduce the influence of some of these processes
(namely, Auger nonradiative scattering, Förster energy transfer, and thermal escape) in
order to focus on the ligand influence on excited state dynamics of CdSe nanocrystals.
Finally, surface state trapping can be studied at low temperature thanks to a broad, sub-
!
92
band gap emission band that appears below 650 nm in CdSe nanocrystals. It is associated
with relaxation of the electron or hole to surface states, followed by emissive, sub-band
gap recombination. The surface states implicated are either selenium
46
or cadmium
vacancies
47
(i.e., Cd or Se dangling bonds, respectively). Historically, CdSe nanocrystals
exhibited sub-band gap emission even at room temperature; progress in surface
passivation has succeeded in reducing this emission as a result of core–shell
32,43
!and/or
ligand engineering.
46,48

Figure 3.12. Red shifts of the first exciton peak for the different R groups in our study, with
increasing red-shifts following the trend tBuSeH < BnSeH < PhSeH.

The optical properties of the ligand exchanged CdSe nanocrystals were investigated by
UV–vis absorption and PL spectroscopies. The samples were dispersed in toluene
(native ligands) or TMU (phenylchalcogenol ligands) and prepared such that they had
similar absorbances. All of the absorbance spectra exhibit a bathochromic shift of the
first exciton peak relative to CdSe(NL) after ligand exchange, with the
phenylchalcogenol ligands causing a red shift following the general trend S < Se ≤ Te
550 570 590 610 630 650
Normalized Absorbance (a.u.)
Wavelength (nm)
NL
PhSeH
BnSeH
tBuSeH
!
93
(Figure 3.8.d). For example, the first exciton peak was red-shifted by 22 meV after
ligand exchange with PhSH, while ligand exchange with either PhSeH or PhTeH both
resulted in a 32 meV red shift relative to CdSe(NL). Bathochromic shifts (<40 meV) in
the optical absorbance spectra of CdSe quantum dots are commonly observed upon
adsorption of thiol ligands.
46
The effect is attributed to the relaxed quantum confinement
resulting from the coupling between the HOMOs of the thiol and the orbitals of the CdSe
valence band.
51,52
Here, the quantum confinement is further relaxed for PhSeH and
PhTeH ligands compared to the PhSH. This enhanced stabilization is related to increased
delocalization of the hole wave function into the Se and Te ligand binding group, as
discussed below. Equivalently, in a finite quantum well picture, the stabilization is due in
part to the size increase observed by TEM (spatial relaxation of the quantum confinement
by increasing the size of the well) and also to the modification of the energetic and
dielectric environment
53
of the nanocrystal (i.e., energetic relaxation of the quantum
confinement by lowering the height of the energy barrier). Interestingly, red shifts of the
first exciton peak were also observed for the different R groups in our study, with
increasing red-shifts following the trend tBuSeH < BnSeH < PhSeH (Figure 3.12.). This
trend likely corresponds to the increasingly proximal aromatic ring to the nanocrystals.
The photoluminescence of CdSe nanocrystals is red-shifted relative to the first excitonic
feature in absorption because of the fine structure of the lowest exciton state.
42

Interestingly, ligand exchange with phenylchalcogenols results in greater red-shifts for
the emission maxima than for the exciton absorbance (Figure 3.13.a); the native ligand
!
94
capped nanocrystals emit at 601 nm, while the PhSH-capped nanocrystals emit at 615
nm, and the PhSeH- and PhTeH-capped nanocrystals emit at 619 nm, corresponding to
shifts by 44 and 57 meV relative to the CdSe(NL) emission, respectively. The greater
shifts observed in emission compared to those observed in absorption reveal that the
magnitude of the exciton fine structure splitting depends on the surface ligands; the dark
exciton (observed in photoluminescence) is more stabilized by ligand exchange than the
bright exciton (observed in absorbance). While the Stokes shift is known to be size
dependent,
54
the expected size effect is much smaller than that observed here. The
exciton fine structure splitting depends on the electron and hole wave function overlap
and on the dielectric surroundings of the nanocrystals,
42
and therefore is expected to be
ligand dependent.
Not surprisingly, the PL intensity is also greatly affected by the ligand exchange. The
steady-state photoluminescence spectra of CdSe nanocrystals ligated either with stearate
or exchanged with phenylchalcogenols clearly show that the ligand exchange results in
strong PL quenching for all three phenylchalcogenol ligands (Figure 3.13.a). Control
experiments performed with neat PhSeH vs PhSeH and Ph
2
PH show comparable
quenching in the absence of Ph
2
PH, suggesting that the phosphine does not play a role in
the observed PL quenching (Figure 3.14). Additionally, control experiments run to
determine the effect of ligand concentration on the PL quenching show that the PL
quenching was not dependent on the concentration of added ligand in the saturated
regime (i.e., 0.018 mmol Ph
2
Se
2
/mg CdSe) used in this study (Figure 3.15.). While this
!
95

Figure 3.13. Steady-state PL spectra of CdSe nanocrystals suspended in toluene (native ligands)
or TMU (phenylchalcogenol ligands). (a) Room temperature steady-state PL spectra of
CdSe(NL) nanocrystals before (green data) and after ligand exchange with PhSH (red data),
PhSeH (blue data), and PhTeH (purple data) with l
ex
= 550 nm. The PL intensities have been
normalized by the absorbances at the excitation wavelength to account for slight differences in
concentration (absorbances = 0.1–0.2 OD). The inset shows a zoom of the phenychalcogenol-
exchanged nanocrystals. (b) Low temperature (77 K) steady-state PL spectra of CdSe(NL)
nanocrystals before (green data) and after ligand exchange with PhSH (red data), PhSeH (blue
data), and PhTeH (purple data) with l
ex
= 550 nm. As a result of the intense PL signals at 77 K,
dilute suspensions were used (absorbances = 2 × 10
–2
OD).

experimental evidence points to the conclusion that ligand coverage may not be the
determining factor in the PL quenching, it is still difficult to definitively rule out this
effect. Thiol ligands are well-known to introduce midgap hole traps in CdSe
nanocrystals, resulting in strong quenching of the photoluminescence.
31,32,46,52,53
The
thiol ligands attract the hole and reduce its wave function overlap with the excitonic
electron counterpart, thereby reducing the coupling between electron and hole and their
!
96
emissive recombination. Interestingly, the PL quenching is even stronger in the case of
PhSeH and PhTeH ligands. This observation is consistent with our DFT calculations of
the phenylchalcogenol ligand energy levels (vacuum, basis B3LYP-LACVP**, Table
3.2.). These calculations suggest that PhSeH and PhTeH possess higher lying HOMO
levels than the PhSH. They can therefore also act as excitonic hole traps, possibly with a
larger driving force for hole transfer than the PhSH ligands. The species created by hole
transfer to the phenylchalcogenol ligands are not emissive in the detected energy range,
as shown by the absence of PL signal below the band gap energy of the nanocrystals and
down to 1.55 eV (800 nm).

Figure 3.14. Control experiments performed to determine the effect of phosphine on the PL
intensity. (a) Room temperature steady-state PL spectra of CdSe(NL) nanocrystals before (green
data) and after ligand exchange with neat PhSeH (red data), and with neat PhSeH and Ph
2
PH
(blue data), with λ
ex
= 550 nm. Both reactions were run under identical conditions as stated in the
experimental. (b) UV-vis spectra of CdSe(NL) nanocrystals before (green data) and after ligand
exchange with neat PhSeH (red data), and with neat PhSeH and Ph
2
PH (blue data). (c) UV-Vis
spectra comparing the reaction run with neat PhSeH (red data) to that run with Ph
2
Se
2
and Ph
2
PH
(purple data).
!
97

Figure 3.15. Plots showing the dependence of PL quenching on the concentration of added
Ph
2
Se
2
and Ph
2
PH.

Table 3.2. HOMO and LUMO levels of phenyl chalcogenols, calculated in the gas phase with
the TITAN software, with B3LYPfunctional and LACVP** basis set, after refining the geometry
at the semi-empirical PM3 model.
HOMO (eV) LUMO (eV)
PhSH -5.95 -0.16
PhSeH -5.88 -0.19
PhTeH -5.65 -0.87

The mechanism of PL quenching via hole transfer to the chalcogenol ligands can be
further investigated by time-resolved PL measurements. At low temperature in
particular, the influence of phonons is reduced, leading to a clearer view of the hole
transfer dynamics. Indeed, time-resolved PL measurements at 77 K show that the band-
edge PL is quenched rapidly in all ligand exchanged nanocrystals (Figure 3.16.b). While
the PL signal of PhSH-capped nanocrystals also exhibit a long-lived component, both
PhSeH- and PhTeH-capped nanocrystals are almost entirely quenched. The characteristic
time for hole quenching is close to the time resolution of the TCSPC apparatus for all
!
98
phenylchalcogenol ligands; in the case of PhSH, the decay can be fitted with a 40 ps
component, while in the case of PhSeH and PhTeH, the fast characteristic time for hole
transfer is smaller than the time resolution (30 ps, Table 3.3). At longer times, all
samples exhibit a slow decay component with a characteristic lifetime on the order of 1
ns, similar to what is observed in the CdSe(NL) sample; however, this component is of

Figure 3.16. Time-resolved PL spectra of CdSe nanocrystals suspended in toluene (native
ligands) or TMU (exchanged ligands). (a) Room temperature time-resolved PL spectra of
CdSe(NL) nanocrystals before (green data, l
em
= 605 nm) and after ligand exchange with PhSH
(red data, l
em
= 614 nm), PhSeH (blue data, l
em
= 619 nm), and PhTeH (purple data, l
em
= 620 nm)
with l
ex
= 550 nm. (b) Low temperature (77 K) time-resolved PL spectra of CdSe(NL)
nanocrystals before (green data, l
em
= 590 nm) and after ligand exchange with PhSH (red data, l
em
= 592 nm), PhSeH (blue data, l
em
= 591 nm) and PhTeH (purple data, l
em
= 601 nm). All signals
have been normalized to their maximum intensity. The black curves show the fits to the data.
The TCSPC apparatus instrument response function (IRF) is given in grey.

!
99

Table 3.3. Summary of fit constants for the PL decays. The experimental decays were fitted by
a convolution of the IRF and of a sum of exponential and stretched exponential decays (with
constants τ
i
and t
st,i
, b
st,i
, respectively) with the Fluofit software.
77 K Room temperature
NL τ
st
= 0.87 ± 0.03 ns,
β
st
= 0.61 ± 0.02
τ
st
= 0.95 ± 0.04 ns,
β
st
= 0.50 ± 0.02
τ
2
= 0.060 ns ± 0.001
PhSH τ
st
= 0.77 ± 0.04 ns,
β
st
= 0.60 ± 0.03
τ
st,2
= 0.042 ± 0.001 ns,
β
st,2
= 0.66 ± 0.01
τ
st
= 0.69 ± 0.007 ns,
β
st
= 0.43 ± 0.002
τ
st,2
< 0.030 ns (= 0.001 ±
0.0003 ns, β
st,2
= 0.48 ±
0.04)
τ
st,3
= 0.036 ± 0.002 ns,
β
st,3
= 0.50 ± 0.02
PhSeH τ
st
= 0.88 ± 0.1 ns,
β
st
= 0.54 ± 0.06
τ
st,2
= 0.023 ± 0.0005 ns,
β
st,2
= 1 ± 0.06
τ
st
= 1.20 ± 0.01 ns,
β
st
= 0.51 ± 0.002
τ
st,2
< 0.030 ns (= 0.001 ±
0.0004 ns, β
st,2
= 0.56 ±
0.04)
τ
st,3
= 0.038 ± 0.003 ns,
β
st,3
= 0.52 ± 0.02
PhTeH τ
st
= 0.89 ± 0.1 ns,
β
st
= 0.48 ± 0.05
τ
st,2
= 0.023 ± 0.0004 ns,
β
st,2
= 1 ± 0.05
τ
st
= 0.87 ± 0.01 ns,
β
st
= 0.42 ± 0.002
τ
st,2
< 0.030 ns (= 0.002 ±
0.0002 ns, β
st,2
= 0.49 ±
0.02)
τ
st,3
= 0.034 ± 0.002 ns,
β
st,3
= 0.48 ± 0.01

!
100
much smaller relative amplitude for PhSeH and PhTeH than for PhSH and NL capped
nanocrystals. At room temperature, all phenylchalcogenol-exchanged CdSe nanocrystals
exhibit a similar, instrument-limited decay component (Figure 3.16.a) that arises from a
combination of hole transfer to the chalcogenol ligand and of phonon-assisted
nonradiative decay pathways. In effect, phonon-assisted relaxation can be assessed by
comparing the PL decay dynamics of the nanocrystals capped with native ligands at room
temperature (Figure 3.16.a) and at 77 K (Figure 3.16.b). At room temperature, the PL is
rapidly quenched with a characteristic time around 60 ps (Table 3.3). This fast
component is not present in the 77 K PL decay, so it is attributed to phonon-assisted
relaxation. The quenching for nanocrystals with PhSH ligands is faster at room
temperature than at 77 K, suggesting that hole transfer to the ligands could be thermally
activated.
Finally, all the nanocrystals presented in this study exhibit a long-lived sub-band gap
emission at 77 K (lifetime determined to be longer than 100 µs from the inverse repetition
rate of the pulse LED illumination source used for these measurements; data not shown)
attributed to surface state trapping (Figure 3.13.b).
46,47
Interestingly, the exchange with
PhSH resulted in an increase of the surface state emission intensity relative to the native
ligands. This observation is consistent with the literature. Baker and Kamat attribute the
increase of surface state emission to selenium vacancies; the exchange with thiols is
thought to result in strong Cd–S bonds, which weaken the surface Cd–Se bonds and
facilitate Se oxidation and increase the number of selenium vacancies and the intensity of
!
101
surface state emission.
46
Here, in the case of the PhSeH and PhTeH, the surface state
emission intensity is reduced relative to the native ligands. This set of observations can
be explained by different scenarios and we think that experimental evidence is lacking to
strictly attribute surface state trapping to electron or to hole trapping, i.e., to Cd dangling
bonds or Se dangling bonds, respectively. Both cases could lead to our observations. In
the case of electron trapping, the increase in surface state emission intensity after thiol
exchange could arise from surface Se oxidation, as in the work of Baker and Kamat—an
effect that would be smaller for Se and Te ligands due to weaker binding strengths with
surface Cd atoms. In the case of surface hole trapping, the highly effective hole transfer
to the PhSeH and PhTeH ligands, discussed above, could compete with the surface state
hole trapping, while the less efficient hole transfer to PhSH would not.
3.4. Experimental
3.4.1. General Considerations
The following chemicals were used as received without further purification. CdCO
3

(99.998% metals basis, “Puratronic” grade, Alfa Aesar), selenium (200 mesh powder,
99.999% metals basis, Alfa Aesar), tri-n-octylphosphine oxide (TOPO, 98%, Alfa Aesar),
tri-n-octylphosphine (TOP, ≥97%, Strem), stearic acid (95%, Sigma-Aldrich),
diphenyldisulfide (Ph
2
S
2
, 98%, TCI America), diphenyldiselenide (Ph
2
Se
2
, 98%, Strem),
diphenylditelluride (Ph
2
Te
2
, 98%, Aldrich), dibenzyldisulfide (Bn
2
S
2
, 98+%, Alfa Aesar),
dibenzyldiselenide (Bn
2
Se
2
, 95%, Alfa Aesar), diphenylphosphine (Ph
2
PH, 98%, Sigma-
!
102
Aldrich). Di-tert-butyldiselenide (
t
Bu
2
Se
2
) was synthesized using a literature procedure.
55

Di-tert-butylditelluride (
t
Bu
2
Te
2
) was also synthesized using a literature procedure.
56

Tetramethylurea (TMU, 99%, Alfa Aesar) was distilled at atmospheric pressure under
nitrogen before use, discarding ∼5–10% residue in the distillation flask. Deuterated
NMR solvents were purchased from Cambridge Isotopes Laboratories and used as
received.
Ligand Exchange Procedure. All ligand exchanges were performed using the following
procedure in 23 mL vials under nitrogen. R
2
E
2
(0.30 mmol) was added to 17 mg of
CdSe(NL) nanocrystals dispersed in 1 mL of toluene. This solution was purged with
nitrogen for 5 min while stirring, followed by the addition of Ph
2
PH (171 mg, 0.920
mmol). Regardless of the R
2
E
2
precursor used, the reaction was allowed to stir for 24 h.
The exchanged CdSe nanocrystals were purified by flocculation with pentane (6 mL),
followed by centrifugation (6000 rpm, 1 min). After each washing step, the particles were
redispersed in TMU (0.5 mL). This procedure was repeated four times to ensure removal
of all free ligands.
Material Characterization. TGA measurements were made on a TA Instruments TGA
Q50 instrument, using sample sizes of ∼5 mg in an alumina crucible under a flowing
nitrogen atmosphere. TGA samples were prepared by fully drying the colloid under
flowing nitrogen at 100 °C for 30 min prior to analysis. FT-IR spectra were recorded on a
Bruker Vertex 80. To obtain quantitative data, an internal standard (Fe
4
[Fe(CN)
6
]
3
, Alfa
!
103
Aesar) was used according to a known procedure.
25
UV–vis spectra were acquired on a
Shimadzu UV-1800 spectrophotometer using a quartz cuvette. TEM was performed on a
JEOL JEM-2100 microscope at an operating voltage of 200 kV, equipped with a Gatan
Orius CCD camera. Samples were prepared from dilute purified dispersions and
deposited onto 300 mesh Formvar-coated copper grids (Ted Pella, Inc.).
1
H,
31
P and
77
Se
NMR spectra were obtained on a Varian 500 spectrometer (500 MHz in
1
H, 202 MHz in
31
P, 95 MHz in
77
Se) with chemical shifts reported in units of ppm. All
1
H chemical shifts
are referenced to the residual
1
H solvent (relative to TMS). EDX spectra were collected
on a JEOL JSM-6610 scanning electron microscope operating at 10 kV equipped with an
EDAX Apollo silicon drift detector (SDD). Multiple regions of a sample deposited on an
aluminum stub were analyzed. PL spectra were collected on a Horiba Jobin Yvon
Nanolog spectrofluorimeter system equipped with a 450 W Xe lamp as the excitation
source and a photomultiplier tube as the detector. The excitation wavelength was 550 nm
for all measurements. The samples were placed in a standard 1-cm path length quartz
cuvette for the room temperature measurements. For low temperature (77 K) PL
measurements, a liquid nitrogen-cooled Dewar was used and the samples were placed in
an NMR tube.
PL Lifetime Studies. TCSPC measurements (30 ps time resolution) were performed using
a R3809U-50 Hamamatsu PMT with a B&H SPC-630 module. Samples were excited
with the 550 nm output of an OPA, pumped by a Ti:sapphire regenerative amplifier
operating at 250 kHz (Coherent RegA 9050). Laser pulses were focused into the sample
!
104
using a 15 cm focal lens, with illumination powers between 3.2 mW and 1.5 µW.
Emission was collected at the emission peak of the nanocrystals at magic angle
polarization relative to the 550 nm excitation. The emission was collimated and then
focused into a monochromator with a 10 cm lens. The monochromator was a CVI
CMSP112 double spectrograph with a 1/8 m total path length in negative dispersive
mode with a 600 g mm
–1
grating blazed at 600 nm. The slit widths were 1.2 mm,
resulting in a 20 nm resolution.
Samples were dispersed in toluene (for the CdSe(NL) nanocrystals) or TMU (for the
phenylchalcogenol-exchanged CdSe nanocrystals) to obtain optical densities between 0.1
and 0.2 OD at the excitation wavelength; cuvettes or NMR tubes were used for room
temperature and 77 K measurements, respectively, as explained above. The excitation
fluences, between 40 µJ cm
–2
and 20 nJ cm
–2
, were chosen depending on the strength of
the photoluminescence signal (i.e., depending on the ligands and on the temperature), so
that excitation resulted in less than one excited electron–hole pair per nanocrystal.
52,53
Attempts were made to measure the lifetime of the longer wavelength sub-band gap
emission of the nanocrystals at 77 K by using a IBH Fluorocube instrument equipped
with a 405 nm LED excitation source at a 10 kHz repetition rate. Although signal was
detected, the extremely long-lived sample luminescence decayed more slowly than the
LED source repetition rate making a quantitative determination of lifetime impossible.

!
105
3.5. Conclusions
In summary, band-edge emission in phenylchalcogenol-exchanged quantum dots is a
competition between (i) radiative relaxation; (ii) hole transfer to the ligands (at 77 K, 40
ps for PhSH, < 30 ps for PhSeH and PhTeH); (iii) surface state trapping, detectable via a
broad, long-lived sub-band gap emission at 77 K; and (iv) phonon-assisted nonradiative
decay pathways at room temperature (∼60 ps). The PhSeH and PhTeH ligands appear to
be more efficient hole acceptors than PhSH ligands. This finding shows that, contrary to
intuition, adding Se atoms to the quantum dots by ligand engineering does not simply
make the nanocrystals “larger” from the PL decay point of view. Mixed coverage of the
surface of nanocrystals by ligands with different end groups (alkyl chains from native
ligands and phenyl groups from exchanged ligands in this case) leads to inhomogeneous
Se or Te shell growth and creates very effective hole-transfer pathways. In particular, we
suggest that hole transfer processes could effectively compete with hole surface state
trapping (followed by emission of a photon) for selenol and tellurol ligands. Finally,
while hole transfer to thiol ligands reduces the band-edge PL lifetime, it has been shown
to increase the ground state depopulation lifetime in transient absorption measurements
of thiol-capped CdSe nanocrystals.
50
Therefore, in the context of solar energy conversion,
hole transfer to the ligand could be a way to delay energy relaxation; efficient hole
quenchers such as selenol or tellurol ligands could become the first step to extract energy
from the system, through a cascade of charge transfers whereby the holes are spatially
pulled away from electrons.
!
106
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Y. Temperature-Dependent Photoluminescence of CdSe-Core CdS/CdZnS/ZnS-
Multishell Quantum Dots. J. Phys. Chem. C 2009, 113, 13545-13550.
(44) Valerini, D.; Cretí, A.; Lomascolo, M.; Manna, L.; Cingolani, R.; Anni, M.
Temperature Dependence of the Photoluminescence Properties of Colloidal
CdSe/ZnS Core/Shell Quantum Dots Embedded in a Polystyrene Matrix. Phys.
Rev. B 2005, 71, 235409.
(45) Morello, G.; De Giorgi, M.; Kudera, S.; Manna, L.; Cingolani, R.; Anni, M.
Temperature and Size Dependence of Nonradiative Relaxation and Exciton-
Phonon Coupling in Colloidal CdTe Quantum Dots. J. Phys. Chem. C 2007, 111,
5846-5849.
(46) Baker, D. R.; Kamat, P. V. Tuning the Emission of CdSe Quantum Dots by
Controlled Trap Enhancement. Langmuir 2010, 26, 11272.
(47) Underwood, D. F.; Kippeny, T.; Rosenthal, S. J. Ultrafast Carrier Dynamics in
CdSe Nanocrystals Determined by Femtosecond Fluorescence Upconversion
!
111
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Meijerink, A. Temperature Antiquenching of the Luminescence from Capped
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the Exciton Luminescence and Decay Kinetics of CdTe and CdSe Quantum Dots.
J. Phys. Chem. B 2004, 108, 17393-17397.
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CdSe Nanoparticles Monitored with Femtosecond Time-Resolution from the
Visible to the IR: Assignment of the Transient Features by Carrier Quenching. J.
Phys. Chem. B 2001, 105, 12286-12292.
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Coupled Quantum Dot-Ligand Systems. J. Phys. Chem. Lett. 2013, 4, 634-640.
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! 112
Chapter 4. Ligand Exchange of Colloidal CdSe Nanocrystals with Stibanates
Derived from Sb
2
S
3
Dissolved in a Thiol-Amine Mixture*
*Published in Chem. Mater. 2014, 26, 6311-6317.

4.1. Abstract
Molecular stibanates derived from the dissolution of bulk Sb
2
S
3
in a binary
ethylenediamine and mercaptoethanol solvent mixture have been studied as capping
ligands for colloidal CdSe nanocrystals. A phase transfer ligand exchange strategy was
utilized to effectively install the stibanate ligands onto the CdSe nanocrystals to form
stable colloidal suspensions in polar solvents, such as formamide. This methodology was
very effective in the removal of insulating native ligands on the as-prepared nanocrystals,
with the resulting stibanate-capped CdSe nanocrystals giving low organic content thin
films upon spin coating with improved interparticle coupling after heating to
temperatures < 300 ˚C. Photoelectrochemical measurements on stibinate-capped CdSe
nanocrystal films showed that this novel ligand leads to a > 25-fold increase in
photocurrent response relative to as-prepared CdSe nanocrystal films.
4.2. Introduction
The performance of colloidal semiconductor nanocrystals in optical and electronic
devices is reaching the tipping point of being technologically relevant. Of particular
importance, colloidal nanocrystals hold promise for the inexpensive manufacture of
semiconducting thin films through solution processing, such as spray coating and roll-to-
! 113
roll printing.
1-3
Focused efforts to make semiconductor nanocrystals attractive for optical
and electronic applications have resulted in impressive research progress in nanocrystal-
based thin films,
4
whereby significant increases in conductivities and carrier mobilities
have been realized in nanocrystal-based thin films.
5
Many of these improvements can be
directly attributed to the development of new ligand exchange strategies that replace the
electrically insulating, long-chain surface ligands (e.g., oleate, stearate, or
alkylphosphonates) from the as-synthesized nanocrystals with smaller ligands that allow
for stronger interparticle coupling.
Our group has had much success using small organic ligands such as tert-butylthiol, as
well as a variety of other small chalcogenol ligands, which very effectively displace the
long-chain ligands present on as-prepared CdSe nanocrystals and allow for improved
interparticle coupling in nanocrystal films.
6,7
Of the ligands studied, the tert-butylthiol
ligand most notably lead to significantly improved photocurrent, carrier mobility, and
power conversion efficiencies in nanocrystal-based devices relative to the as-prepared
and pyridine-exchanged CdSe nanocrystals.
6,8,9
Another approach that has been
employed to enhance the interparticle coupling of nanocrystal-based thin films is to
perform ligand exchange reactions with small inorganic ligands. The newest class of
inorganic ligands for semiconducting nanocrystals is made up of various halide,
pseudohalide and halometallate ions.
10-13
Use of these ligands has allowed for efficient
electronic passivation and significant improvements in the charge transport of
nanocrystal-based devices; for example, iodide-capped CdSe nanocrystals give high
! 114
carrier mobilities of 12 cm
2
V
–1
s
–1
.
11
Electron mobilities as high as 25 cm
2
V
–1
s
–1
were
observed in chloride-capped CdSe nanocrystal-based field-effect transistors (FETs);
however, the nanocrystals were sintered, which may contribute to the high carrier
mobilities.
13
Earlier successes were also achieved with thiocyanate
14,15
and sulfide
ligands,
16,17
both of which were shown to facilitate strong interparticle coupling and
charge transport between nanocrystals once assembled in thin films. Lastly,
chalcogenidometalates are another early and well-known class of inorganic ligands that
have been utilized for semiconductor nanocrystals.
18-21
While originally hydrazinium-
based (e.g., (N
2
H
5
)
4
Sn
2
S
6
, (N
2
H
4
)
2
(N
2
H
5
)
2
In
2
Se
4
), these ligands can now be prepared and
utilized as non-hydrazinium-based salts (e.g., (NH
4
)
4
Sn
2
S
6
, K
4
SnTe
4
). Since their first
use by Talapin and coworkers, these strongly binding ligands have proven their utility by
allowing for improved conductivity in nanocrystal thin films.
18-20
Many of the chalcogenidometalates utilized by Talapin and co-workers were originally
prepared by Mitzi at IBM for the solution-processing of semiconductor thin films from
dissolved semiconductor inks.
22-24
Their general synthesis involves the simple
dissolution of bulk metal chalcogenides in hydrazine with stoichiometric chalcogen
through a process known as dimensional reduction. In an analogous way, we recently
developed an alternative route to dissolve semiconductors using a binary solvent mixture
comprised of ethylenediamine (en) and a thiol without the need for added chalcogen.
This solvent mixture possesses the remarkable solvent power to rapidly dissolve a variety
of notoriously insoluble semiconductors at room temperature and ambient pressure.
25-27

! 115
Intrigued by the parallels between hydrazine and our en/thiol solvent system, we set out
to investigate the molecular nature of dissolved Sb
2
S
3
(herein used as a model system) in
en/mercaptoethanol (ME), and subsequently installed these dissolved species as ligands
on CdSe nanocrystals. Indeed, these molecular stibanate species can be successfully
utilized for the rapid and efficient ligand exchange of native stearate ligands present on
as-prepared CdSe nanocrystals to produce colloidally stable, stibanate-capped CdSe
nanocrystals. Solution-processed nanocrystal thin films of the stibanate-capped CdSe
nanocrystals demonstrate improved interparticle coupling relative to the as-prepared
nancorystals, which leads to markedly higher electrochemical photocurrent generation.
4.3. Results and Discussion
4.3.1. Characterization of the Dissolved Sb
2
S
3
Species
Stibanates from dissolved Sb
2
S
3
were chosen as the model chalcogenidosemimetalate
ligands in this study because of their promise for electronic and optoelectronic
applications.
21,28-30
In the bulk form, Sb
2
S
3
possesses a direct band gap (E
g
= 1.5-1.8 eV)
with a strong absorption coefficient (7.5 × 10
4
cm
–1
at 550 nm).
31
Small
chalcogenidosemimetalate clusters of Sb
2
S
3
prepared by the dissolution of the bulk
material in hydrazine and excess chalcogen have previously been utilized as inorganic
ligands on CdSe nanocrystals.
21
While small chalcogenidometalate ligands are generally
known to enhance the performance of nanocrystalline thin film field-effect transistors by
increasing their conductivity by several orders of magnitude when compared to organic
! 116
ligands, not much is known about the properties of stibanate-capped CdSe
nanocrystals.
21,29
Herein, the stibanate ligands were prepared by dissolution of bulk
Sb
2
S
3
in a 1:40 vol/vol mixture of mercaptoethanol (ME) and ethylenediamine (en)
(Figure 4.1.). Due to the novelty of this solvent system, characterization of the dissolved
Sb
2
S
3
species was needed in order to gain insights into the possible identity of the
resulting chalcogenidosemimetalate ligands. Chemical information on the stibanate
species was obtained using a combination of ICP-AES elemental analysis,
thermogravimetric analysis (TGA), ESI-MS, Raman and FT-IR spectroscopies.

Figure 4.1. Photograph before (left) and after (right) dissolution of 20 mg bulk Sb
2
S
3
powder in
2 mL en and 0.05 mL ME. After dissolution the light-yellow solution is optically transparent and
free of visible scattering.
Initial information on the composition of the dissolved Sb
2
S
3
was obtained by analyzing
the chemical composition of the dissolved species after it had been evaporated to dryness.
The chemical composition of resulting sample was confirmed by ICP-AES and
combustion elemental analysis to have the formula C
4.6
H
12.0
N
2.1
SbS
2.3
(Table 4.1.). This
composition, made up of roughly 47 wt% organics (C, H, and N) and 53 wt% heavier
inorganic elements (Sb and S), is corroborated by TGA on the same material, which
! 117
demonstrates a 45% mass loss up to 500 ºC. Further analysis of dissolved Sb
2
S
3
in the
binary en/ME solvent mixture implies that rather than a singular species, the stibanates
likely exist as multiple molecular species in solution. A negative-ion ESI-MS spectrum
acquired from a dilute solution of Sb
2
S
3
in en/ME (240 µg mL
–1
) indicates the presence
of four primary stibanate species that can be classified into two groups containing either
one or two Sb atoms (Figure 4.2.). The main ion clusters observed at m/z 228.9 and m/z
272.9 are consistent with the formulas [SbS
2
C
2
H
4
O]
–
and [SbS
2
C
4
H
8
O
2
]
–
, respectively.
Assignment of the other two ion clusters at m/z 416.8 and m/z 460.8 are consistent with
the formulas [Sb
2
S
4
C
2
H
5
O]
–
and [Sb
2
S
4
C
4
H
9
O
2
]
–
, respectively.
Table 4.1. Raw data for elemental analysis on the stibanate ligands. C, H, N, S and Sb are
reported as a mass fraction (wt/wt%).
Element Method Amount used Result
C Combustion analysis, CHNS/O
Analyzer
1.967 mg
2.136 mg
11.85%
11.91%
H Combustion analysis, CHNS/O
Analyzer
1.967 mg
2.136 mg
2.57%
2.57%
N Combustion analysis, CHNS/O
Analyzer
1.967 mg
2.136 mg
6.23%
6.21%
S Combustion analysis, Sulfur
Determinations Using the
LECO SC-432DR
4.176 mg 15.67%
Sb ICP-AES 9.60 mg 26.1%

! 118

Figure 4.2. A negative-ion ESI-MS spectrum of Sb
2
S
3
dissolved in 1:40 vol/vol mixture of
en/ME at a concentration of 240 µg mL
–1
.
The possible identity of the stibanate species present in the en/ME solution can be
garnered from these elemental analysis and ESI-MS data. ESI-MS clearly indicates that
the dissolved Sb
2
S
3
produces a mixture of compounds that contain ME ligands. ICP-AES
elemental analysis suggests the presence of en because of the additional C and N content.
From these data a few formulae are possible, including, [SbS][C
2
H
4
SO][C
2
H
9
N
2
],
[Sb][C
4
H
8
S
2
O
2
][C
2
H
9
N
2
], [Sb
2
S
3
][C
2
H
4
SO][C
2
H
9
N
2
] and [Sb
2
S
2
][C
4
H
9
S
2
O
2
][C
2
H
9
N
2
].
Characterization of similar species in the literature may point to the structure of some of
the stibanate species present here (Figure 4.3). The reaction between Sb(III)
isopropoxide and ME is known to yield a compound [Sb][C
4
H
8
S
2
O
2
(H)] similar in
formula to the one found in this work (i.e., [Sb][C
4
H
8
S
2
O
2
][C
2
H
9
N
2
]).
32
The structure of
[Sb][C
4
H
8
S
2
O
2
(H)] is 3-coordinate about Sb with approximate pyramidal geometry. One
of the two ME ligands forms a five membered chelate ring with Sb in
! 119
[Sb][C
4
H
8
S
2
O
2
(H)], while the other ligand is bonded only through the S atom.
32
FT-IR
spectroscopic analysis on Sb
2
S
3
dissolved in en/ME after it had been evaporated to
dryness displayed similar features to that reported for [Sb][C
4
H
8
S
2
O
2
(H)], while also
displaying additional spectral features originating from en (Figure 4.4.). For example, the
presence of bands at 1279 cm
–1
, 1045 cm
–1
, and 543 cm
–1
can be matched to the ν(S-C),
ν(C-O), and ν(Sb-OC) bands from [Sb][C
4
H
8
S
2
O
2
(H)], respectively. The broad ν(N-H)
band around 3300 cm
–1
likely results from protonated en counterions in the species
investigated here.
33
Finally, the stretching band corresponding to the ν(S-H) thiol group
at 2560 cm
–1
is not present in the spectrum of the dissolved Sb
2
S
3
in en/ME, thus
indicating the likely deprotonation of the thiol with concomitant thiolate coordination to
the Sb center. Dimeric species, which may be similar to the [Sb
2
S
2
][C
4
H
8
S
2
O
2
][C
2
H
9
N
2
]
and [Sb
2
S
3
][C
2
H
4
SO][C
2
H
9
N
2
] compounds, are also well known in the literature. In fact,
the majority of molecular species produced by dissolving Sb
2
S
3
in caustic sulfidic
solutions are dimeric with two Sb atoms bound together through two bridging sulfur
atoms.
34,35
Interestingly, the Raman spectrum of the stibanates from dissolved Sb
2
S
3
in
en/ME displays two key bands at 300 and 411 cm
–1
in the region characteristic for Sb-S
bonds, and which closely match with the calculated stretching frequencies of the neutral,
dimeric Sb
2
S
2
(SH)
2
complex (Figure 4.5.).
34

! 120

Figure 4.3. Proposed chemical structures of the stibanate ligands derived from Sb
2
S
3
dissolved
in a solvent mixture of en/ME.

Figure 4.4. FT-IR spectrum of the dried stibanate ligands derived from Sb
2
S
3
dissolved in a
solvent mixture of en/ME.
! 121

Figure 4.5. Raman spectrum of a solution of Sb
2
S
3
dissolved in en/ME. Bands at 300 and 400
cm
–1
appear after the dissolution of Sb
2
S
3
, while the band at 337 cm
–1
is due to the (CN) torsion
from en in the liquid state.
4.3.2 Ligand Exchange Using Stibanates from Dissolved Sb
2
S
3
Due to its well-known synthesis and surface chemistry, CdSe was used as the model
nanocrystal platform to study the ligand exchange chemistry described herein. The CdSe
nanocrystals were synthesized using a previously reported procedure whereby tri-n-
octylphosphine selenide (TOPSe) is injected into a hot solution of Cd(stearate)
2
in TOPO
and stearic acid.
36
The resulting nanocrystals displayed a first excitonic transition at
~593 nm in the UV-vis absorption spectrum, indicating a CdSe nanocrystal diameter of
ca. 4.4 nm.
37
Transmission electron microscopy (TEM) revealed that the CdSe
nanocrystals possess a spherical morphology with a diameter of 4.9 ± 0.5 nm, which is in
general agreement with empirical sizing by UV-vis spectroscopy. TGA data showed a
single mass loss event of 28% up to 500 ºC. This is consistent with previous work,
6,7
and
! 122
is indicative of a native ligand shell comprised mainly of anionically bound stearate,
which makes the as-synthesized CdSe nanocrystals readily dispersible in non-polar
solvents such as toluene.

Figure 4.6. (a) Photograph before (left) and after (right) phase transfer of CdSe nanocrystals
from the nonpolar toluene phase to the polar formamide phase caused by exchange of the native
aliphatic ligands with stibanate ligands. (b) ζ-potential of stibanate-capped CdSe nanocrystals in
formamide with a photograph of the colloidal suspension given as an inset. (c) TGA traces
comparing the mass loss events observed for the stibanate ligands, as-prepared CdSe
nanocrystals, and ligand exchanged CdSe nanocrystals. (d) FT-IR spectra of as-prepared and
ligand exchanged CdSe nanocrystals after drying. The spectra are off-set and normalized in
intensity to the 2089 cm
–1
ν(C≡N) stretch of a measured Fe
4
[Fe(CN)
6
]
3
internal standard (not
shown).

! 123
Facile removal of the stearate native ligands through installation of a stibanate ligand
shell was accomplished in a single step using a simple phase transfer procedure. In a
typical two-phase system, a solution of as-prepared CdSe nanocrystals in toluene was
stirred with a dilute solution of Sb
2
S
3
/en/ME in formamide (Figure 4.6.a). Complete
transfer of the CdSe nanocrystals from the top toluene layer to the bottom formamide
layer was achieved within 5 min, indicating facile removal of the insulating stearate
native ligands. Subsequent extraction of the top toluene layer followed by a two-step
purification process yielded a stable colloidal suspension of stibanate-capped CdSe
nanocrystals in formamide. Following ligand exchange, the stibanate-capped CdSe
nanocrystals could be easily filtered through a 0.1 µm filter and were colloidally stable in
formamide for periods > 1 month. Zeta potential measurements on the stibanate-capped
CdSe nanocrystals corroborate the empirically observed colloidal stability and revealed a
ζ-potential of -32 mV (Figure 4.6.b), indicating (i) a stable suspension and (ii) that
negatively charged surface ligands are responsible for the electrostatic stabilization of
these particles in formamide. In an effort to show the general scope of this ligand
exchange procedure, additional ligands and nanocrystalline cores were screened. The
en/ME solvent system and ligand exchange procedure was successfully applied to other
chalcogenidosemimetalate/chalcogenidometalate ligands on CdSe nanocrystals, as well
as using the stibanate ligands on other nanocrystalline cores (Figure 4.7.). Successful
dissolution and ligand exchange was achieved with the following semiconductors, as
evidenced by efficient phase transfer: As
2
S
3
, As
2
Se
3
, Sb
2
Se
3
, SnS, and ZnS on CdSe
! 124
nanocrystals. Furthermore, ligand exchange with the stibanates was amenable to other
inorganic nanocrystals, such as CdS/CdSe core/shell nanocrystals and Pt nanocrystals,
which were synthesized according to literature procedures.
38,39
Similar to the CdSe
nanocrystals, the UV-vis spectra of the CdS/CdSe core shell nanocrystals before and after
ligand exchange with the stibanate ligands show no apparent scattering or blue shift of
the first exciton peak, indicating that the nanocrystals are colloidally stable and not
etched, respectively.

Figure 4.7. (a) The generality of this approach is illustrated by performing the ligand exchange
using a variety of dissolved bulk semiconductors on CdSe nanocrystals. The same procedure was
used in all cases with Sb
2
S
3
being replaced with the semiconductors listed in the figure. (b) The
scope of this ligand exchange procedure was further illustrated by performing the ligand
exchange on various inorganic nanocrystals using the stibanate ligands. (c) Solution absorption
spectra of the CdSe nanocrystals before and after ligand exchange with various dissolved bulk
semiconductors. (d) Solution absorption spectra of CdS/CdSe core/shell nanocrystals before and
after ligand exchange with stibanate ligands.

! 125
Control reactions were carried out to gain insight into the ligand exchange chemistry and
colloidal stability of the CdSe nanocrystals. Formamide proved to be irreplaceable for
the phase transfer reaction. When other solvents, such as water or N,N-
dimethylformamide (DMF), were used as the second phase during the phase transfer, the
CdSe nanocrystals did not fully transfer from the toluene phase or were not colloidally
stable, respectively. The importance of formamide is possibly due to its high dielectric
constant (ε = 111) and donor number (DN = 39.8), which allow for the electrolytic
dissociation of cations, while also facilitating strong electrostatic stabilization of the
stibanate-capped CdSe nanocrystals.
10,40
Control experiments performed in the absence
of dissolved Sb
2
S
3
(i.e., using just en/ME/formamide as the polar phase) underwent
complete phase transfer; however, the resulting nanocrystals were not colloidally stable
upon purification, indicating the importance of the stibanate ligands for proper
stabilization of the particles in formamide.
4.3.3. Efficacy of Ligand Exchange
The efficacy of the phase-transfer ligand exchange was probed using a number of
complementary techniques. First, the ligand exchange was investigated using TGA. All
TGA samples were extensively dried and purified, so as to specifically study the surface
bound ligands, before running the analysis under flowing nitrogen (5 ºC min
–1
to 100 ºC,
then 10 ºC min
–1
to 500 ºC, Figure 4.6.c). The data showed that the as-prepared CdSe
nanocrystals possess a single high temperature mass loss of 28% beginning at 300 ºC,
which was assigned to the loss of stearate native ligands (vide supra). After ligand
! 126
exchange with the stibanate ligands, the high temperature mass loss corresponding to the
native ligands was no longer observed by TGA, indicating quantitative removal of the
native ligands. Additionally, in the stibanate-capped CdSe nanocrystal sample, new mass
loss events are observed that are consistent with those observed for dried Sb
2
S
3
/en/ME
(i.e., the stibanate ligands). Both traces display three mass loss events occurring in the
same temperature ranges (100-200 ºC, 200-300 ºC, and 300-500 ºC). Comparison of the
mass% ratios of these three steps relative to the total mass loss gives 66:27:7 and
61:29:10 for the stibanate ligands and stibanate-capped CdSe nanocrystals, respectively.
Therefore, it is likely that the new ligand shell on the ligand exchanged CdSe
nanocrystals is composed of roughly the same species obtained after drying the
Sb
2
S
3
/en/ME solution. Finally, it should be noted that the decomposable ligand mass
went from 28% to 3.8% after ligand exchange with the stibanate ligands, which is in
agreement with the 5 wt% Sb coverage achieved after ligand exchange (vide infra). The
low degree of organic content, in addition to the low temperature (< 300 ºC) recovery of
the pure inorganic material, leads to the ability to make nanocrystal films from this ink
that are not prone to micro- and macroscopic cracking.
Further validation of the ligand exchange efficacy was obtained by FT-IR spectroscopy
using samples dispersed in KBr (Figure 4.6.d). Samples were dried at 100 ºC before
analysis, and the spectra were normalized in intensity to the 2089 cm
–1
ν(C≡N) stretching
frequency of a Fe
4
[Fe(CN)
6
]
3
internal standard (see Experimental Methods). The use of
an internal standard allowed for normalization of the observed FT-IR bands; however,
! 127
quantification could not be performed due to extensive overlap of the ν(C-H) stretching
bands before and after ligand exchange. Regardless, it was clear that the strong features
associated with the ν(C-H) stretching bands of the native stearate ligands (at 2800-3000
cm
–1
) were greatly reduced after ligand exchange with the stibanate ligands. Stearate
ligand removal was also verified using
13
C NMR spectroscopy by comparing the spectra
of the as-prepared and ligand–exchanged nanocrystals after a typical acid digestion and
extraction procedure using aqua regia.
21
The
13
C NMR spectra in benzene-d
6
of the
liberated ligands suggest clearly shows that after ligand exchange there are no carbon
resonances characteristic of stearate ligands at 178.3, 33.8, 32.4, 30.2, 25.0, 23.2 and 14.4
ppm (Figure 4.8.). In accordance with the TGA, ESI-MS, and elemental analysis data, it
is likely that the low intensity bands present in the ν(C-H) stretching region after ligand
exchange result from the presence of en and ME moieties of the stibanate ligands.
Moreover, there is a clear band in the range of 3000-3600 cm
–1
likely resulting from a
combination of ν(N-H) and ν(O-H) bands, which indicate the presence of en and/or ME
species present after ligand exchange.
33
In accordance with TGA, complete removal of
the organics present after ligand exchange was achieved after annealing to 300 ºC, as
evidenced by the loss of intensity in the ν(C-H) stretching region (vide infra).

! 128

Figure 4.8.
13
C NMR spectra of the acid digestion products from the as-prepared (top) and
ligand-exchanged (bottom) CdSe nanocrystals in benzene-d
6
.
While TGA and FT-IR provide information on the organic ligands present before and
after ligand exchange, elemental analysis and XPS can provide quantitative information
on the stibanate ligands present after ligand exchange. Elemental analysis by ICP-AES
of the stibanate-capped CdSe nanocrystals revealed a Cd:Sb atomic ratio of 10:1,
indicating the adsorption of stibanate ligands on the surface of the CdSe nanocrystals
(Table 4.2.). The Cd:Se ratios were also analyzed using ICP-AES elemental analysis to
confirm that the integrity of the CdSe nanocrystal core was maintained after ligand
exchange. Elemental analysis showed that the Cd:Se ratios remained generally
unchanged after ligand exchange (i.e., Cd:Se for as-prepared and stibanate-capped CdSe
nanocrystals were 1.07:1 and 1.11:1, respectively).

! 129
Table 4.2. ICP-AES analysis of as-prepared and ligand exchanged CdSe nanocrystals. Cd, Se,
and Sb reported as a mass fraction (wt/wt%).
Sample Amount used Cd Se Sb
as-prepared 15.42 mg 54.9% 36.2% N/A
ligand exchanged 11.29 mg 50.9% 32.1% 5.05%

The CdSe nanocrystal surface chemistry was further probed by XPS analysis.
Information on the newly installed stibanate ligands was obtained from a high-resolution
scan of the Sb3d spectral region (Figure 4.9.a). In this region there were only two Sb3d
peaks that displayed the typical splitting pattern for Sb atoms in Sb
2
S
3
,

with the Sb3d
5/2

and Sb3d
3/2
peaks measured at 530.1 eV and 539.4 eV, respectively.
41
The peaks are also
consistent with those obtained from XPS analysis of the stibanates from dissolved Sb
2
S
3

in en/ME, which possess binding energies of 529.8 eV and 539.1 eV for the Sb3d
5/2
and
Sb3d
3/2
peaks, respectively. Furthermore, the lack of shoulder peaks in the high-
resolution Sb3d spectrum confirms that there is no elemental Sb or Sb
5+
species present,
and that the binding energy of the stibanates present on the surface of the CdSe
nanocrystals is consistent with Sb in the 3+ oxidation state.
After ligand exchange, the CdSe nanocrystals possessed a mean diameter of 5.1 ± 0.6 nm,
which is not statistically different from the as-prepared nanocrystals. Further analysis of
the TEM micrographs shows that the CdSe nanocrystals also retain their morphology
after phase transfer and ligand exchange (Figure 4.10.). Additionally, the interparticle
! 130
distance is greatly decreased after ligand exchange (0.1-0.4 nm) when compared to the
as-prepared CdSe nanocrystals (2-4 nm), indicating that the stibanate ligands may

Figure 4.9. (a) High-resolution XPS spectrum of the Sb3d region of the stibinate-capped CdSe
nanocrystals after ligand exchange. (b) Solution absorption and PL spectra of as-prepared and
ligand exchanged CdSe nanocrystals.
enhance the interparticle coupling between nanocrystals. Solution UV-vis absorption
spectra of the ligand exchanged nanocrystals support the assertion that the ligand
exchange procedure does not affect the integrity of the CdSe nanocrystal cores. This is
supported by the similarity between solution absorption spectra for the as-prepared and
stibanate-capped CdSe nanocrystals (Figure 4.9.b). Not only are the excitonic features
maintained after ligand exchange, but there are no signs of nanocrystal etching or gross
aggregation. Solution photoluminescence spectra were also taken before and after ligand
exchange, with the band-edge PL intensity being greatly quenched after ligand exchange
(Figure 4.9.b). This quenching is characteristic for chalcogenolate and
chalcogenidometalate ligands on CdSe,
7,20
which offers further evidence of ligand
binding to the nanocrystal surface.
! 131

Figure 4.10. TEM micrographs of the as-prepared and ligand exchanged CdSe nanocrystals.
4.3.4. Nanocrystal Film Characterization
The photoelectrochemical properties of stibanate-capped CdSe nanocrystal thin films
were investigated to determine the utility of this material for future optoelectronic
applications. Thin films of both the as-prepared and the ligand-exchanged nanocrystals
were deposited onto ITO coated glass by spin coating to produce films of roughly the
same optical density. Photocurrent measurements were performed using a three-
electrode electrochemical cell, with two 472 nm LEDs for illumination (Figure 4.11.a,
4.12.). The ligand-exchanged CdSe nanocrystal films give repeatable, high anodic
photocurrent densities under illumination (ca. -45 µA cm
–2
at 125 mV relative to Pt
pseudoreference). On the other hand, the as-prepared CdSe nanocrystal films displayed
very small photocurrents (< 2 µA cm
–2
at 125 mV relative to Pt pseudoreference). It is
likely that the increased photoelectrochemical response from the stibanate-capped CdSe
nanocrystals is a consequence of removing the insulating native ligands, which are
known to impede conduction. UV-vis spectra of the heat treated CdSe nanocrystal films
! 132
shows the as-prepared CdSe nanocrystal films remain quantum confined with a clearly
defined first exciton peak, while the ligand-exchanged CdSe nanocrystal films display a
broadening and red-shifting of the first excitonic peak that is likely a result of enhanced
interparticle coupling (Figure 4.11.b).
15
Quantitative FT-IR on the nanoparticles before
and after heat treatment was used to determine how the ligand exchange and heat
treatment affected the loss of organic ligands. The intensities for the as-prepared CdSe
nanocrystals before and after 150 °C heat treatment showed no difference, while heat-
treatment to 300 ºC on the ligand-exchanged CdSe nanocrystal films results in nearly
complete removal of the organics (Figure 4.13.), corroborating our hypothesis that the
observed increase in photocurrent response after ligand exchange is due to insulating
ligand loss. Moreover, X-ray diffraction analysis of the stibanate-capped CdSe
nanocrystal films before and after heating to 300 ºC did not reveal any appreciable
change in diffraction peak breadth (Figure 4.14.); the absence of diffraction peak
sharpening suggests that nanocrystal sintering was not in significant effect.

! 133

Figure 4.11. (a) Photocurrent response for ligand-exchanged CdSe nanocrystal films heat treated
to 300 °C and as-prepared CdSe films heat treated to 150 oC during a potential scan from −300 to
+300 mV relative to a Pt pseudoreference electrode with chopped 472 nm illumination. (b) UV-
vis spectra of spin-cast nanocrystal films of as-prepared and ligand exchanged CdSe after mild
annealing (to 150 °C for as-prepared nanocrystals and 300 °C for ligand exchanged nanocrystals).

Figure 4.12. Photocurrent response for ligand exchanged CdSe nanocrystal films heat treated to
300 °C (bottom) and as-prepared CdSe films heat-treated to 150 ºC (top). Potentials are relative
to a Pt pseudoreference electrode, and chopped illumination was performed with 472 nm LEDs.
! 134

Figure 4.13. FT-IR spectra of the dried, as-prepared and dried, ligand exchanged CdSe
nanocrystals dispersed in KBr, before and after heat treatment. The spectra were normalized to
the 2089 cm
-1
ν(C≡N) stretching peak of Fe
4
[Fe(CN)
6
]
3
internal standard.

Figure 4.14. XRD patterns comparing the ligand exchanged CdSe nanocrystals before and after
heating to 300 ºC.
! 135
4.4. Experimental
4.4.1. Materials and Methods
CdCO
3
(99.998% metals basis, “Puratronic” grade), selenium (200 mesh powder,
99.999% metals basis), tri-n-octylphosphine oxide (TOPO, 98%), sulfur (precipitated,
99.5%) and formamide (99.5+%) were purchased from Alfa Aesar. Tri-n-octylphosphine
(TOP, ≥ 97%) was purchased from Strem. Stearic acid (95%), ethylenediamine (en, ≥
99%), 2-mercaptoethanol (ME, 99+%), and Sb
2
S
3
(99.995%) were purchased from
Sigma-Aldrich. Na
2
S hydrate (> 60% Na
2
S) was purchased from Lancaster. All reagents
and solvents were used as received without further purification.
Dissolution of bulk Sb
2
S
3
. Prior to ligand exchange, a Sb
2
S
3
/ethyelendiamine
(en)/mercaptoethanol (ME) stock solution was prepared by mixing Sb
2
S
3
(20 mg) with en
(2.00 mL) and ME (0.05 mL) at room temperature under a nitrogen atmosphere. The
solution was stirred until complete dissolution of the bulk Sb
2
S
3
occurred (ca. 30 min) to
yield a transparent light-yellow solution that was then handled in air.
Ligand exchange. For ligand exchange of the as-prepared CdSe nanocrystals with the
dissolved Sb
2
S
3
, a phase transfer procedure was utilized. In this procedure, a
Sb
2
S
3
/en/ME stock solution (0.2 mL) was diluted with formamide (2.0 mL) under a
nitrogen atmosphere. A second immiscible suspension of the as-prepared CdSe
nanocrystals in toluene (5.5 mL, 2.5 mg mL
–1
) was layered on top of the polar phase.
This resulted in immediate phase transfer of the CdSe nanocrystals from the top toluene
! 136
phase to the bottom polar phase. After 5 min, quantitative phase transfer was achieved as
evidenced by the colorless toluene phase and colored polar phase. The top toluene layer
was then fully extracted. The stibanate-capped CdSe nanocrystals suspended in the polar
formamide phase were then purified by flocculation with 12 mL of acetone,
centrifugation (6000 rpm, 1 min), with the supernatant being discarded. This purification
procedure was repeated 2 × with final redispersion in 1 mL of formamide. To test for
agglomeration, the final product was centrifuged (6000 rpm, 1 min) and no solid was
deposited, indicating minimal or no agglomeration.

Material characterization. The chemical composition of the CdSe nanocrystals was
determined using inductively coupled plasma atomic emission spectroscopy (ICP-AES)
(Galbraith Labs, Knoxville, TN). Additionally, carbon, hydrogen, and nitrogen elemental
analysis was made using a PerkinElmer2400 Series II CHNS/O Analyzer (Galbraith
Labs, Knoxville, TN). Samples (25 mg) were prepared for elemental analysis by drying
under vacuum at 90 ºC for 24 h. Electrospray ionization mass spectroscopy (ESI-MS)
was performed using a Waters LCT premier operated using electrospray ionization with
negative ion detection. FT-IR spectra were recorded on a Bruker Vertex 80 spectrometer.
To obtain quantitative data, a weighed internal standard (Fe
4
[Fe(CN)
6
]
3
, Alfa Aesar) was
used according to a known procedure.
6
The resulting spectra were then baseline
corrected using a rubber band mode with 5 iterations and normalized to the internal
standard peaks from 2160-2020 cm
–1
. All other samples were dropcast onto ZnSe
windows and dried at 100 ºC. Raman spectra of the solutions were recorded under
! 137
ambient conditions using a Horiba Jobin Yvon spectrometer, equipped with a liquid
sample holder. An excitation source of 785 nm from a diode laser was employed at a
power level of 50 mW. Thermogravimetric analysis (TGA) measurements were made on
a TA Instruments TGA Q50 instrument, using sample sizes of ~5-10 mg in an alumina
crucible under a flowing nitrogen atmosphere. TGA samples were prepared by fully
drying the samples under flowing nitrogen at 100 °C for 30 min prior to analysis.
Transmission electron microscopy (TEM) images were obtained using a JEOL
JEM2100F (JEOL Ltd.) microscope operating at 80 kV. Samples for TEM studies were
prepared by drop-casting a stable suspension of nanocrystals in toluene or formamide on
a 400 mesh Cu grid coated with a lacey carbon film (Ted Pella, Inc.). Average particle
diameters and standard deviations were derived by measuring a minimum of 200
individual particles per experiment and by averaging over multiple images. X-ray
photoelectron spectroscopy (XPS) was obtained on a Kratos Axis Ultra X-ray
photoelectron spectrometer with the 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. UV-vis spectra were acquired on a
Shimadzu UV-1800 spectrophotometer using a quartz cuvette. Photoluminescence (PL)
spectra were collected on a Horiba Jobin Yvon Nanolog spectrofluorimeter system
equipped with a 450 W Xe lamp as the excitation source and a photomultiplier tube as the
! 138
detector. The excitation wavelength was 550 nm for all measurements. ζ-potential
measurements were obtained using a Zetasizer Nano-ZS (Malvern Instruments, U.K.).
Suspensions were measured in formamide after being transferred into disposable
capillary cells (DTS1070, Malvern). ζ-potential was calculated from electrophoretic
mobility using Henry’s equation in the Smoluchowski limit.
42
Film deposition and photoelectrochemistry. Spin coating was carried out using a Laurell
Technologies Corporation WS400Ez- 6NPP-LITE Single Wafer Spin Processor, fitted
with a continuous nitrogen purge and a heat lamp to assist with film drying. The as-
prepared CdSe nanocrystals were spun at room temperature from a 20 mg mL
–1

suspension in toluene at 1000 rpm for 60 s (acceleration = 770 rpm s
–1
), while the ligand-
exchange nanocrystals were spun at 80-100 ºC from a 20 mg mL
–1
suspension in
formamide at 1000 rpm for 60 s (acceleration = 770 rpm s
–1
). When pre-cleaned
microscope slides were employed as a substrate, no preparation was necessary. For ITO-
coated glass (Delta Technologies 7 × 50 × 0.7 mm
3
, R
s
= 5-15 Ω), the substrates were
sequentially sonicated in isopropyl alcohol for 15 min and acetone for 15 min before
being ozone treated for 20 min. Thermal treatment was carried out on a thermostatically
controlled Al heating block under flowing nitrogen. Photoelectrochemical response was
performed using a BASi Epsilon-EC potentiostat. A quartz cuvette was used with a Pt-
wire counter electrode and a Pt-wire pseudo-reference electrode. The working electrode
was a CdSe nanocrystal film on ITO-coated glass. An aqueous 0.050 M Na
2
S/0.005 M
sulfur electrolyte was made from nitrogen-sparged deionized water, and care was taken to
! 139
avoid exposure of the electrolyte solution to oxygen throughout the course of the
experiments. For photoelectrochemical experiments, two 472 nm LEDs mounted ca. 6
cm from either side of the sample were used to illuminate the working electrode. The
illumination intensity was measured using a Newport model 818-ST photodiode
connected to a Newport 2832-C dual-channel power meter. Measurements were taken
for both illumination directions and summed, giving a result of 4.9 mW cm
−2
. The total
illumination area of the CdSe working electrode was 2.3 cm
2
.
4.5. Conclusions
In summary, have developed a new route for ligand exchange of CdSe nanocrystals by
using molecular stibanates derived from the dissolution of bulk Sb
2
S
3
in an en/ME
solvent system. Using a simple phase transfer method, we are able to enact efficient
ligand exchange on the as-prepared CdSe nanocrystals; that is, rapid and complete
replacement of the insulating native ligands for stibnate ligands is achieved in a matter of
minutes. The resulting stibinate-capped CdSe nanocrystals are colloidally stable in polar
solvents, which allows for facile solution processing into nanocrystal thin films. The
stibanate-capped CdSe nanocrystals in the resulting thin films show a bathochromic shift
in the first exciton peak after heating to 300 ºC without significant sintering and grain
growth, suggestive of improved interparticle coupling from removal of the insulating
native ligands. This, in turn, leads to a pronounced improvement in photoelectrochemical
response of the nanocrystal films relative to those derived from as-prepared nanocrystals.
We believe this unique solvent system can be optimized for other combinations of
! 140
chalcogenidometalate ligands and semiconductor nanocrystals, thereby further enriching
the toolbox of surface modification techniques for nanocrystals.
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!
! 144
Chapter 5. Dissolution of Sn, SnO and SnS in a Thiol–Amine Mixture:
Insights into the Identity of the Sn-Based Molecular Precursor for Solution
Processed SnS

5.1. Abstract
Alkanethiol / ethylenediamine – based solvent systems have the ability to readily dissolve
semiconductors under ambient conditions thus enabling facile solution processing of
electronic semiconductors. Despite numerous publications utilizing this solvent system,
there is no direct information regarding the chemical identity of the dissolved species.
Herein we examine the molecular solute formed after dissolution of Sn, SnO and SnS in a
solvent mixture comprised of 1,2-ethanedithiol (EDT) and 1,2-ethylenediamine (en).
Using a suite of
119
Sn solution NMR, Raman and FT-IR spectroscopies,
bisethanedithiolate tin(II) was identified as the single tin species present in all three (Sn,
SnS and SnO) / en / EDT solutions, despite the various bulk precursor compositions (Sn,
SnO and SnS) and oxidation states (Sn
0
and Sn
2+
). All three inks can be converted to SnS
using a mild annealing step (~350 ºC). Therefore, lower-cost Sn precursors can be used
as raw materials for the solution deposition of crystalline, phase-pure SnS.
5.2. Introduction
Solution processing of inorganic semiconductors for the fabrication of affordable,
scalable electronic thin-films is emerging as a commercially viable alternative to
traditional high vacuum deposition methods.
1
Using this approach, semiconductor thin
! 145
films can be printed onto a various substrates (e.g. plastics, glasses) at low temperatures
using existing high-throughput technologies, such as roll-to-roll manufacturing, inkjet
printing, and spray coating.
2
Innovations in the field of molecular (semiconductor) inks
have lead to high quality electronic thin-films, which have shown utility in
photovoltaics,
1,3
thermoelectrics,
4
thin-film transistors,
5
and phase change memory
materials.
6
Molecular semiconductor inks, as opposed to colloidal nanocrystalline inks, are
comprised of molecular solutes in solution and are typically prepared using one of two
methods. The inorganic compound can be introduced into solution as a premade soluble
precursor,
7,8
or through the direct dissolution of bulk materials having the desired
elements for the recovered semiconductor.
9-13
Direct dissolution offers a number of
advantages over the former of the two methods; in a single step, cheap bulk materials can
be used as elemental sources for high quality semiconductors using a simple “dissolve
and recover” process. Additionally, this approach offers enhanced flexibility in the
semiconductor’s composition because the stoichiometry of the final product is often
dependent on the stoichiometry of the dissolved bulk material or combination of thereof.
The challenge of this approach lies in the fact that most bulk metals and semiconductors
are insoluble in common solvents, as dissolution from bulk to molecular form involves
breaking the strong covalent bonds that build up these frameworks.
Hydrazine is one such solvent capable of dismantling semiconducting and/or metallic
frameworks to yield high-quality soluble inks. Hydrazine in the presence of elemental
! 146
chalcogen was first employed by Mitzi and coworkers at IBM as an effective solvent for
the dissolution of bulk metal chalcogenides.
9-11
Dissolution using hydrazine is facilitated
by reactive chalcogen E
2-
(E
2−
=S
2−
, Se
2−
, and Te
2−
) species formed via the in-situ
reduction of chalcogen. Reaction between E
2-
and bulk metal chalcogenides yields
soluble hydrazinium chalcogenidometallate species, which can be easily deposited to
form high-quality metal chalcogenide thin-films after annealing. A number of hydrazine-
processed semiconductors have demonstrated impressive device efficiencies on par with
those manufactured using vacuum-deposition techniques.
14
For instance, hydrazine-
processed CuIn
x
Ga
(1-x)
Se
2
(CIGS) photovoltaic devices have shown record power
conversion efficiencies (> 15%).
1

The foundation for these impressive device performances lies in having a clear
understanding of the dissolution chemistry and the relevant precursor complexes in
hydrazine.
15
A number of studies have investigated the structure of the dissolved
precursor species, the solute/solvent interactions and the molecular condensation
reactions that occur upon annealing and semiconductor recovery.
11,16-19
This knowledge
has provided key insights into the phase formation mechanism of the as-deposited films
and ultimately allowed for more control over the resulting crystalline phase, composition,
electronic band gap and film morphology. Therefore, it is critical to understand the
dissolution chemistry for the purpose of device performance optimization.
In addition to hydrazine, alternative solvent systems comprised of alkanethiols and 1,2-
ethylenediamine (en) have been developed by our group for the facile dissolution of a
! 147
wide array of metals and semiconductors.
12
The dissolved materials can be directly
recovered as high quality thin-films (e.g. Se, Te, SnS and various V
2
VI
3
chalcogenides),
or converted to the corresponding sulfides using the alkanethiol as a sulfur source (i.e.
bulk oxides to metal sulfide thin-films).
13,20-21
The solvent system can also be used to
synthesize compositionally controlled semiconducting alloys (e.g. Sb
2
Se
3-x
S
x
,
CuIn(SxSe
1–x
)
2
) or even used to mix and match dissolved elements to form
semiconductors having higher order compositions (e.g. Cu
2
ZnSnS(Se)
4
, Sb
2
Se
3
and
SnTe).
22-25
Furthermore, the solvent system has shown commercial promise for the
fabrication of high performance electronics.
26
For example, alkanethiol-en based
CZTSSe photovoltaic devices show efficiencies up to 7.34%.
27
Additionally, proof-of-
principle flexible electronics have been made from deposition of Cu
2
S and Cu
2
Se on
flexible polyimide substrates.
28
While the utility of the solvent system for technologically relevant applications has been
shown, there are no systematic or direct spectroscopic studies aimed at determining the
atomistic details of dissolved inorganic species. As was shown with the hydrazine
system, experimental insights into the solvent/solute chemistry aided device optimization.
Similarly, we believe developing a thorough understanding of the different molecular
species formed upon dissolution of bulk semiconductors in the alkanethiol / en solvent
system is an essential for further developments to be made. A deep chemical
understanding not only gives researchers a finer degree of control over the current
! 148
dissolution, processing and recovery steps but may also open a window of opportunity so
new and innovative applications can emerge.
In this study, we take a look into the alkanethiol-en solution chemistry and determine the
precursor complexes formed after dissolution of three bulk Sn-precursors. Previously,
both SnS and SnO precursors have been dissolved using the alkanethiol-en solvent
system for the deposition of tin monosulfide (SnS);
20-21,24
however, nothing has been
reported on the nature of the dissolved species. By focusing on Sn-based precursors we
were able to use the NMR active
119
Sn isotope for solution NMR studies. Additionally,
upon deposition, these Sn-based inks yield high quality semiconducting SnS, a promising
light-absorbing material for photovoltaic devices. SnS is a narrow band gap
semiconductor (direct E
g
= 1.32 eV and indirect E
g
= 1.08 eV) possessing an
orthorhombic (Pnma) layered crystal structure.
29
Although less prevalent in the literature,
theory predicts that the energy conversion efficiency of single-junction SnS-based solar
cells can reach up to 32%, with existing SnS solar cells reaching up to 4.4% efficiency.
30

Thus, solution processing of SnS could allow for the fabrication of an affordable
photovoltaic made using inexpensive and earth abundant elements.
5.3. Results and Discussion
5.3.1. Dissolution of Sn, SnO and SnS
Solutions of Sn, SnO and SnS were prepared by dissolution of the bulk material, at room
temperature and pressure in a 10:1 vol/vol mixture of en/EDT according to previous
! 149
reports.
12
The solubilities, expressed as wt"% solute in the saturated solution of en/EDT at
room temperature and pressure, of both SnO and SnS have been previously
determined.
20,21
SnO when dissolved at room temperature for a period of 15 h was
soluble up to 15 wt% in a 4:1 vol/vol ratio of en/EDT, while SnS was soluble up to 12
wt% in an a 1:11 vol/vol ratio of en/EDT after mild heating (50 ºC, overnight). For this
paper a lower concentration of 4.5 wt% was used for all three systems due to the lower
solubility limit of the Sn system. At this concentration both SnS and SnO
semiconductors were fully soluble after stirring under flowing N
2
(g) for 24 h. Longer
reaction times were required for the dissolution of Sn; even after ~ca. 5 days a minimal
amount of metallic Sn remained in the vial. Despite the reduced wt % solubility of Sn,
all solutions gave yellow-tinted, optically transparent solutions that were free of visible
scattering and stable for a period of months (Figure 5.1A).

! 150

Figure 5.1. Spectroscopic characterization of Sn, SnO and SnS dissolved in a 10:1 vol/vol 1,2-
ethylenediamine / 1,2-ethanedithiol–solvent mixture. A) Photograph, B)
119
Sn NMR spectra, C)
FT-IR spectra, and D) Raman spectra, of SnO (blue data), SnS (purple data) and Sn (green data)
dissolved in the en / EDT solvent system.
5.3.2. Characterization of the Dissolved Sn Species.
119
Sn NMR spectroscopy was used to probe the structure of the dissolved Sn species in
solution. The
119
Sn isotope in particular posses’ ideal NMR properties such as a spin ½,
making spectra and coupling easy to understand; a natural abundance of 8.58% and a
receptivity 25.7 times higher than
13
C, allowing for reasonable NMR acquisition times;
and a large chemical shift (δ) range of ca. 6500 ppm, making the chemical shift sensitive
to even small structural changes.
31
Furthermore, large volumes of reported experimental
NMR data on tin compounds allow empirical correlations between the NMR parameters
and molecular structure to be made. One such correlation of particular interest to this
study is that between the δ(
119
Sn) range and the Sn coordination number.
32
Irrespective of
! 151
the formal oxidation state, Sn
+2
or Sn
+4
, the Sn coordination number can be inferred from
δ(
119
Sn) regions;

however due to solvent effects, self association and temperature effects
complementary techniques are often required for full characterization.
31

The
119
Sn NMR spectra, obtained directly after dissolution of the Sn, SnO and SnS
precursors in the en/EDT solvent system, reveal a single sharp peak (δ = 217 ppm) within
the range of +1000 ppm to -1000 ppm for all three systems (Figure 5.1B). Thus all three
systems, having various bulk precursor compositions (Sn, SnO and SnS) and oxidation
states (Sn
0
and Sn
2+
), yield a single equivalent Sn compound on the NMR timescale.
Direct identification of the dissolved Sn compound by matching the spectroscopic
signature at 217 ppm to reported
119
Sn chemical shifts proved futile. For example, the
δ(
119
Sn) did not fall in the range of common thiostannate anions (e.g. SnS
4
4-
and SnS
3
2-
)
which appear in a narrow range between 40 and 80 ppm.
33,34
However, the observed
δ(
119
Sn) is consistent with that reported for Sn-tetrathiolate compounds. For instance, the
chemical shift values of simple Sn-tetrathiolates can range from +26 ppm to +138 ppm
for Sn(SBu
t
)
4
and Sn(SEt)
4
, respectively.
35-37
This range is extended further downfield
when tin is bound by a chelating thiolate, take (SCH
2
CH
2
S)
2
Sn for example, which has a
chemical shift of +277 ppm.
31
No obvious Sn-Sn spin-spin coupling also suggests the
dissolved Sn species is mononuclear and contained a single Sn atom. From these initial
NMR spectra it was inferred that regardless of the bulk Sn precursor, a singular four-
coordinate Sn-thiolate compound was formed in solution.
! 152
Raman spectroscopy on the Sn/en/EDT solutions supports the
119
Sn NMR data. This
characterization technique offers a similar opportunity to that of
119
Sn NMR in that is
allows for the study of the vibrational modes of the dissolved Sn species while solvated
in solution. Figure 5.1C shows the Raman spectra of the (Sn, SnO or SnS)/en/EDT
systems. New vibrational modes observed after dissolution of the Sn-precursors are
assigned to the formation of new bonds, assuming that the en/EDT solvent vibrations
remain consistent (Figure 5.2.). Raman spectra obtained after dissolution of Sn, SnS and
SnO in the en/EDT solvent system are essentially identical to each other, with new bands
appearing at 197 cm
−1
, 264 cm
-1
and 322 cm
-1
. These bands fall in the same range as
those reported for the tetrahedral SnS
4
vibrations of spirocyclic (SCH
2
CH
2
S)
2
Sn(IV)
whose most intense bands occur at 190 cm
−1
, 323 cm
-1
and 356 cm
-1
. Discrepancies
between the dissolved Sn data and the literature data may result from deviations in the
complex symmetry, which is well-known to change the number of Sn-S vibrations and
their wavenumbers.
38
It should be further noted there are no significant peaks associated
with the symmetric stretching v(Sn-N) at 178 cm
-1
which is readily observed upon
coordination of amine compounds with spirocyclic (SCH
2
CH
2
S)
2
Sn(IV). On this basis, it
is likely that the en solvent does not bind to Sn, leaving Sn ligated solely by thiolate
ligands.
! 153

Figure 5.2. Raman spectrum of the 10:1 vol/vol en/EDT solvent system.
Removal of the excess en and EDT solvent provided an opportunity to study the
unsolvated Sn compound using an expanded set of suitable characterization techniques.
Excess solvent was removed by drop casting the Sn-inks onto a glass substrate and
heating to 107 ºC under vacuum for 24 hours. FT-IR spectra comparing the three dried
Sn-based inks are shown in Figure 5.1D. In this region, between 500 cm
-1
to 3500 cm
-1
,
bands originate from the organic component of the dried Sn-species, as any bonds to Sn
such as Sn-S or Sn-N appear from 323-342 cm
-1
and 400-440 cm
-1
, respectively.
39
All
three spectra are the same and generally match the FTIR spectra of EDT, with the
exclusion of any bands corresponding to the ν(S–H) thiol group at 2560 cm
–1
, implying
that the EDT ligands are fully deprotonated. This evidence is in good agreement with the
proposed identity of the dissolved Sn compound as a four coordinate Sn with two
ethanedithiolate ligands.
! 154
5.3.3. Control Reactions and Determination of Sn Oxidation State
Direct characterization of the dissolved Sn species gave us a general idea of the chemical
framework however our inability to match the δ(
119
Sn) to reported values for plausible
compounds (i.e. Sn chelated by EDT, en or a mixture thereof) and lack of insight on the
Sn oxidation state necessitated further investigation. In an effort to nail down the
chemical framework and identify the oxidation state of the dissolved Sn species, we
devised a control system where a plausible compound was synthesized and characterized
along side the dissolved Sn species. Given the simplicity of the solvent system (i.e. it
only contains en, EDT and either Sn, SnO or SnS) and the previous NMR data indicating
a four-coordinate Sn-S compound, it was hypothesized that the dissolved species could
be similar in nature to bis(ethanedithiolate)tin(IV) (1). 1 is well-known known in the
literature and was synthesized following a literature procedure where an aqueous mixture
of 1,2-ethanedithiol and potassium hydroxide was added to a solution of SnCl
4
⋅5H
2
O in
water.
40
The resulting white precipitate was recrystallized from DCM to yield color-less
crystals that were confirmed to be 1 by NMR spectroscopy, melting point determination
and electrospray ionization-mass spectrometry (ESI-MS) (vide infra). The chemical
composition of the sample was determined by combustion CHNS elemental analysis,
which revealed a C:S:H:N ratio of 1:1.02:1.02:0 which is consistent with 1.

! 155

Figure 5.3.
119
Sn NMR spectra of 1 in various solvents showing the dependence of δ(
119
Sn) on
Sn coordination number. The green and purple shaded areas indicates the δ(
119
Sn) ranges for four
and six coordinate tin compounds, respectively.
1 has been characterized using NMR with a reported δ(
119
Sn) of 277 ppm, δ(
1
H) of 3.23
ppm and δ(
13
C) of 35.6 ppm in thiolformaldehyde.
41
While this value does not match the
δ(
119
Sn) of the dissolved Sn species it was plausible that the unique properties of the
en/EDT solvent system could cause shifts in the δ(
119
Sn) value reported in
thioformaldehyde. To test this hypothesis, 1 was characterized in several solvents using
119
Sn NMR. The δ(
119
Sn) of 1 varied quite drastically depending on the solvent between
+500 ppm and -500 ppm; recall that the chemical shift regions are characteristic of tin
coordination environments and a chemical shift in the region form 20 to 280 ppm
indicates a four coordinate tin thiolate while upfield chemical shift range from -125 to
! 156
-515 ppm indicates six-coordinate tin (Figure 5.3.).
42
In CDCl
3
the dissolved compound
gave a single peak in each NMR spectra with a δ(
119
Sn) of 281 ppm, δ(
1
H) of 3.22 ppm
and δ(
13
C) of 35.9 ppm, all of which are in good agreement with literature values (Figure
5.4.).
40
In pure EDT, the δ(
119
Sn) shifted slightly to 269 ppm and was significantly
broadened, implying dynamic exchange between the ethanedithiolate ligands and the
EDT solvent. Deprotonation of EDT resulted in the binding of an additional
ethanedithiolate ligand to 1 rather than exchange with the existing ligands, as evidenced
from the a large downfield shift to -352 ppm, a region characteristic for six-coordinate Sn
compounds.
43,44
Similarly, when en was used as a solvent it appeared to coordinate to 1
yielding a six coordinate compound having a δ(
119
Sn) at -263 ppm. The negative ion
ESI-MS spectrum of 1 in the presence of en reveals that 1 is readily chelate by en, as
evidenced by the main ion cluster observed at m/z 362.9, which corresponds to the
formula [SnS
4
C
4
H
8
N
2
C
2
H
7
]
-
(Figure 5.5.). Finally, when 1 was dissolved in the en/EDT
solvent system the δ(
119
Sn) appeared in the six-coordinate region (-346 ppm) and was
significantly broadened. Raman spectra of the system reveals a strong band at 178cm
-1

which is assigned to both the observed δ(N-Sn-S) and ν(Sn-N) bands for the 1 ligated by
amine ligands. Therefore 1, when dissolved in the en/EDT solvent, is likely coordinated
by en (Figure 5.6.).
! 157

Figure 5.4.
1
H (top) and
13
C (bottom) NMR spectra of 1 in CDCl
3
.

! 158

Figure 5.5. Negative-ion ESI-MS spectrum of 1 dissolved in CHCl
3
in the presence of en. The
main ion cluster observed at m/z 362.9 is consistent with the formula [SnS
4
C
4
H
8
N
2
C
2
H
7
]
-
.

Figure 5.6. Raman spectrum of 1 dissolved in en/EDT. The band at 178 cm
-1
is assigned to
both the δ(N-Sn-S) and ν(Sn-N).
! 159
From the above results it is clear that the dissolved species was not 1. As a Sn(IV)
compound, 1 is easily chelated by ethanedithiolate or en, and thus readily forms six-
coordinate compounds in their presence, while the dissolved Sn compound remained
four-coordinate in the en/EDT solvent system. Four coordinate Sn(IV) compounds have
the propensity to form compounds with a coordination number of six in solution, whereas
the analogous four coordinate Sn(II) compounds do not (Figure 5.7.).
45
The coordination
geometries of Sn(IV) and Sn(II) compounds are well known: four coordinate Sn(IV) is
tetrahedral, has a vacant 5d orbital and can accept further ligands to take on an octahedral
environment; four coordinate Sn(II) has seesaw geometry due to the presence of a
nonbonding electron pair and no longer posses an empty orbital, as the 4pz orbital is
coordinated by donor ligand.
31,46
Based on the coordination behavior of the dissolved Sn,
it appeared that the dissolved Sn species was in the Sn
+2
oxidation state, rather than the
Sn
+4
oxidation state.

Figure 5.7. Sn(II) vs Sn(IV) coordination geometry.

! 160
The influence of Sn oxidation state (Sn
+2
vs Sn
+4
) on δ(
119
Sn) is rarely reported in the
literature, thus further control experiments were required to identify the dissolved Sn
species. As a simple test, Sn
2+
and Sn
4+
in the form of

SnCl
2
⋅2H
2
O and SnCl
4
⋅5H
2
O,
respectively, were dissolved in the en/EDT solvent system and characterized using
119
Sn
NMR (Figure 5.8.). Dissolution of SnCl
4
⋅5H
2
O yielded a single upfield resonance at
-347 ppm, consistent with our previous experiment showing Sn
4+
, from BEDT-Sn(IV),
forms a six coordinate compound when dissolved in en/EDT. While dissolution of
SnCl
2
⋅2H
2
O yielded a single resonance nearly identical to that of the dissolved Sn, SnO
or SnS in the four coordinate region at 210 ppm. It is likely that this peak represents the
same compound as that formed from the dissolution of Sn, SnO and SnS, as a mixture of
SnCl
2
⋅2H
2
O + SnO yielded a single peak at 216.6 ppm. This NMR data further validates
the hypothesis that the dissolved Sn species is in the +2 oxidation state and remains four
coordinate in the en/EDT solvent system.

! 161

Figure 5.8.
119
Sn NMR spectra of SnCl2⋅2H
2
O (top) and SnCl
4
⋅5H
2
O (bottom) dissolved in the
en/EDT solvent system.
5.3.4. SnS Recovery Upon Annealing
Coupled thermogravimetric analysis / differential scanning calorimetry – mass
spectroscopy (TGA/DSC-MS) was used to better understand the ligand composition and
decomposition pathways of 1 so it could function as a benchmark for the characterization
of the dried Sn species (Figure 5.9.a). This technique simultaneously collects TGA and
DSC data in one measurement, thereby differentiating exothermic and endothermic
events associated with (degradation) or without mass loss, while also determining the
identity of the gases evolved during the heating process by mass spectroscopy. Thermal
analysis (50-500 °C, 10 °C min
–1
, under argon flow) of 1, showed a single step mass loss
event at ca. 300 ºC where 30.1% of the total mass was lost. By DSC, three main
transitions were observed. An endotherm not associated with a mass loss event was seen
at 186 ºC which corresponded to the melting point as determined by melting apparatus
! 162
175-183 ºC and the literature reported value for 1 ranging between 180 – 183 ºC.
40
This
was followed by an exotherm (264 ºC, 10.9% mass loss) and endotherm (314 ºC, 19.2%
mass loss) coincident with the major mass loss event observed in TGA. The
decomposition products observed in the mass spectra at both 264 ºC and 314 ºC show
similar products and are consistent with the pyrolysis of ethanedithiolate ligand
decomposition (Figure 5.10.).

Figure 5.9. Thermogravimetric, spectroscopic and optical characterization of the recovered SnS.
A) TGA/DSC traces comparing the decomposition of 1 to the dried Sn, SnO and SnS / en / EDT
systems. B) FT-IR spectra of the SnO-ink dried to 107 °C for 24 hours under vac and after
annealing to 350 °C. C) Powder XRD patterns of the recovered SnS obtained after annealing the
dried Sn-inks to 350 °C under flowing N
2
(g). D) Tauc plots showing the indirect band gaps of the
recovered SnS.
! 163

Figure 5.10. Mass spectra of the evolved gases obtained from TGA/DSC at ~315 ºC for all
samples.
The general TGA/DSC-MS data for 1 shows a striking resemblance to that observed for
the three dissolved and dried Sn, SnO, SnS systems. All of the TGA traces display the
same thermal decomposition route from 100 ºC to 500 ºC which takes place in a single
step at ca 300 ºC and results in a mass loss ranging from 41.7-46.3%. If that mass at 500
°C is approximated as pure SnS, it reveals that the dissolved Sn species typically
contained an average of 44.5% organic matter after thoroughly drying the samples at 107
ºC under vac. This organic composition is close to that calculated for 1 (for
Sn(IV)S
4
C
4
H
8
, the mass percent of SnS is 49.8%, leading to an organic composition of
50.2%) and points to the presence of two EDT ligands per Sn atom, as in 1. By DSC, the
dried samples display similar characteristics to 1 with exotherms at ~275 ºC and
endotherms at ~315 ºC, associated with an average of 15% mass loss and ~30% mass
loss, respectively. Similarly, the major volatile decomposition products observed in the
mass spectrum at ~275 ºC and ~315 ºC are the same as those seem during the
0.0E+00
2.0E-10
4.0E-10
6.0E-10
8.0E-10
1.0E-09
1.2E-09
1.4E-09
1.6E-09
1.8E-09
2.0E-09
10 30 50 70 90 110 130 150
Ion Current (A
-1
)
m/z
BEDT-Sn
SnO
Sn
SnS
C 3 H 7 S 2 (M
+
)
120
M
+
-15
105
M
+
-56
64
M
+
-60
60
M
+
-75
45
M
+
-93
27
Ar
+

39
Ar
++
20
! 164
decomposition of 1, with no signal from ethylenediamine being observed at m/z = 30.
Therefore it is reasonable to infer that EDT ligates Sn in a nature similar to 1 with en
being completely expelled during the drying stage of the sample preparation (∼107 °C
under vacuum, 24 h).
Conversion of the colorless Sn-inks to the corresponding black SnS semiconductor was
achieved using a mild annealing step. The minimum annealing temperature required for
recovery of pure SnS was determined using TGA, which showed that organic thermal
decomposition step occurring at ca. 300 ºC was completed by 350ºC. FT-IR of the dried
Sn-inks lacked any features attributed to organics and confirmed that 350 ºC was an
appropriate annealing temperature for recovery of the inorganic material (Figure 5.9.B).
Powder X-ray Diffraction (XRD) confirmed that the material obtained after annealing the
dried Sn inks to 350 ºC under flowing N
2
(g) was crystalline and phase-pure SnS. Figure
5.9.C shows the diffraction patterns of the SnS obtained from each system (Sn, SnO, and
SnS) after annealing at 350 °C. The diffraction patterns for all three systems are
identical, which coincides with the finding that all the dissolved species are the same.
The XRD diffractograms exhibited a 100% intensity peak (2θ = 31.9°), indexed to the
(111) reflection of orthorhombic Pnma SnS (JCPDS no. 01-073-1859) with all other
peaks matching to the SnS Herzenbergite phase. No SnO or SnS impurities were
observed by TGA for any of the samples further indicating their purity. Analysis of 5
randomly selected large areas of powdered samples by scanning-electron microscope
energy-dispersive X-ray spectroscopy (SEM-EDX) gave an average elemental
! 165
composition of 53 atom % Sn and 47 atom % S for Sn, 57 atom % Sn and 43 atom % S
for SnO and 56 atom % Sn and 44 atom % S for SnS recovered after annealing the dried
inks to 350 °C. These atomic percentages matched those obtained for recovered SnS
from previous studies and also for the commercial SnS standard.
20
The optical properties
of the recovered SnS after annealing at 350 °C were investigated by diffuse reflectance
UV–vis–NIR spectroscopy using an integrating sphere (Figure 5.11.). Figure 5.9.D
shows the indirect band gap approximations obtained from the Tauc plots of the UV/vis
diffuse reflectance spectra. All three SnS samples displayed the same approximate
indirect (E
g,ind
= 1.1 eV) band gap transitions, which agree well with both literature
values (vide supra).

Figure 5.11. Diffuse reflectance spectra of SnS obtained from annealing the Sn, SnO and SnS
inks at 350 ºC under N
2
(g).

In order to gain more insight into the recovery mechanism, the two Sn(IV) control
systems (dissolved 1 and dissolved SnCl2⋅2H
2
O (top)) were recovered using mild
annealing of the dried inks to 350 ºC under N
2
(g). Due to the difference in dissolved Sn
! 166
species in these systems compared to the Sn(II) systems, it was hypothesized that these
systems would yield crystalline SnS
2
. However, the XRD patterns of the recovered
material show the annealing step produced phase pure SnS (Figure 5.12.). Additionally,
analysis of 5 randomly selected large areas of powdered samples by SEM-EDX gave an
average elemental composition of 56 atom % Sn and 44 atom % S for 1 and 60 atom %
Sn and 40 atom % S for SnCl
4
⋅5H
2
O recovered after annealing the dried inks to 350 °C.
Thus the Sn(IV) systems yield SnS with slightly higher atomic percentages of Sn than the
SnS obtained from the Sn(II) systems. This was an interesting result because it shows
that a variety of bulk tin precursors in either the 0, +2 or +4 oxidation states can be
dissolved in the en / EDT solvent system to yield either Sn(II) or Sn(IV) dissolved
species which are then easily recovered as phase pure SnS.

Figure 5.12. XRD patterns of the recovered 1 and SnCl
4
⋅5H
2
O controls from the en / EDT
solvent system.
! 167
5.4. Experimental
5.4.1. Materials and Methods
1,2-Ethylenediamine (en, ≥ 99%), tin (II) sulfide (SnS, 99.99%) and tin (IV) chloride
pentahydrate (SnCl
4
⋅5H
2
O, 98%) were purchased from Sigma-Aldrich. 1,2-Ethanedithiol
(EDT, 98+%), tin (Sn, 99.85%), and tin (II) oxide (SnO, 99%) were purchased from Alfa
Aesar. Tin (II) chloride dihydrate (SnCl
2
⋅5H
2
O, 98%) was purchased from Strem
Chemicals. Ethylenediamine and ethanedithiol were purged using nitrogen gas for a
minimum of 10 minutes before use in the dissolution experiments; all other reagents and
solvents were used as received without further purification.
Preparation of Sn-inks. The (Sn, SnO and SnS) / en / EDT solutions were prepared by
mixing the appropriate Sn precursor (50 mg) with nitrogen purged en (1.00 mL) and EDT
(0.1 mL) solvents at room temperature under a nitrogen atmosphere. The solutions were
then stirred for ca. 5 days to yield transparent light-yellow solutions. Complete
dissolution was observed for the SnS and SnO bulk materials. Sn did not fully dissolve;
excess Sn was removed by filtration through a 450 nm PTFE membrane.
Recovery of Products. Solvent removal entailed heating the drop-cast Sn solutions on to a
temperature controlled hot plate to 100 ºC under flowing nitrogen followed by heating to
107 ºC under vacuum (24 h). The Sn solutions were annealed on a temperature-
controlled hot plate to 350 ºC under flowing nitrogen followed by a slow cool down to
room temperature.
! 168
Material characterization.
119
Sn Solution NMR spectra were recorded on a Varian 500
MHz spectrometer at 186.5 MHz. Spectra were obtained at room temperature without
deuterium locking of the main magnetic field. Each experiment used a 10 µs pulse width,
a 0.5 s relaxation delay and 1024 scans. The spectral window from -1000 to +1500 ppm
was used to search for the signals in 500 ppm increments. δ(
119
Sn) were referenced to the
internal standard of the probe. The NMR samples were prepared in 5mm J-young tubes
under N
2
or air.

Electrospray ionization mass spectroscopy (ESI-MS) was performed
using a Waters LCT premier operated using electrospray ionization with negative ion
detection. The carbon, hydrogen, nitrogen and sulfur chemical composition of the dried
solutions was analyzed using Flash 2000 CHNS Elemental Analyzer. Samples (3 mg)
were prepared for elemental analysis by drying under vacuum at 107 ºC for 24 h. Raman
spectra of the solutions were recorded under ambient conditions using a Horiba Jobin
Yvon spectrometer, equipped with a liquid sample holder. An excitation source of 785
nm from a diode laser was employed at a power level of 50 mW. FT-IR spectra were
recorded on a Bruker Vertex 80 spectrometer. All samples were drop cast onto ZnSe
windows and either dried at 107 ºC under vacuum for 24 hours or annealed to 350 ºC
then immediately cooled. The resulting spectra were then baseline corrected using a
rubber band mode with 5 iterations and normalized to the largest peak.

Thermal analysis
measurements were made on a Netzsch DSC/TGA - MS/IR STA449 F1 using sample
sizes of ~20 mg. Argon was used as protective and flow gas. Mass flow controllers were
set to 40 mL/min for the protective gas and 10 mL/min for the purge gas. The thermal
! 169
range examine was 50 – 500
0
C with a heating rate of 10
0
C/min. All samples were
prepared by fully drying the samples under vacuum at 107 °C for 24 hours prior to
analysis. Powder X-ray diffraction (XRD) data was collected with a Rigaku Ultima IV
diffractometer in parallel beam geometry (2 mm beam width) using Cu Kα radiation (λ =
1.54 Å).

Scanning electron microscope-energy dispersive X-ray spectroscopy (SEM-
EDX) was used for elemental analysis on a JEOL JSM-6610 scanning-electron
microscope.

Diffuse reflectance UV-Vis-NIR spectroscopy was performed using the
reflectivity mode in a Perkin-Elmer Lambda 950 that was equipped with a 150 mm
integrating sphere. The samples were prepared by thoroughly grinding 9 mg of SnS with
290 mg of MgO and using a powder sample holder at the end of the integrating sphere.

5.5. Conclusions
In conclusion, we have identified the molecular structure of the dissolved species
produced after dissolution of Sn, SnO and SnS in the en/edt solvent system. Our finding
reveal that upon dissolution of the bulk precursors a single Sn species is produced in
solution whereby Sn(II) is chelated by two ethanedithiol/-ate molecules. This compound
has yet to be reported in the literature due to its apparent instability therefore this solvent
system has the ability to isolate a relatively unstable species. Mild annealing of the Sn
inks yields crystalline, phase pure SnS. Therefore, inexpensive sources of Sn such as
elemental Sn or SnO can be used as a convenient Sn source for the deposition of SnS.

! 170
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Abstract (if available)

Abstract Solution processing of inorganic semiconductors for the fabrication of affordable, scalable electronic thin‐films is emerging as a commercially viable alternative to traditional high vacuum deposition methods. Using this approach, semiconductor thin‐films can be printed onto a various substrates (e.g. plastics, glasses, or textiles) at low temperatures using existing high‐throughput technologies. In this work, we have developed a number of strategies to synthesize, understand and enhance nanocrystalline and molecular inks for the deposition of high‐quality inorganic semiconducting films. ❧ Compositionally controlled SnₓGe₁₋ₓSe nanocrystal inks were synthesized for the first time as an environmentally sound alternative to lead‐ and cadmium‐based chalcogenides. The band gaps, lattice parameters, and morphologies of the alloys were easily tuned via nanocrystal composition. Interestingly by alloying we showed facile band gap tuning, which is otherwise difficult to achieve in layer structure IV–VI semiconductors. ❧ As synthesized, semiconductor nanocrystal inks are not ready for use in nanocrystal‐based devices due the presence of electrically insulating long‐chain organic surface ligands. While these native ligands ensure chemical and colloidal stability, they must be removed in order to enhance the conductivity of the nanocrystalline inks for electrically conductive films. With this in mind, we have developed a number of ligand exchange techniques whereby the electronically insulating ligands on cadmium selenide (CdSe) nanocrystals are replaced with smaller molecules to enhance the electron mobility of the resulting thin‐films. In one approach we installed small chalcogenol ligands via the in situ reduction of R₂E₂ (R = ᵗBu, Bn, Ph

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University of Southern California Dissertations and Theses

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Asset Metadata

Creator Buckley, Jannise J. (author)

Core Title Semiconductor inks for solution processed electronic thin‐films

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 02/17/2016

Defense Date 11/16/2015

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

Tag inorganic semiconductors,molecular inks,nanocrystal surface functionalization,nanocrystal synthesis,OAI-PMH Harvest,solution‐processing,thin‐films

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Brutchey, Richard (committee chair), Dapkus, Paul Daniel (committee member), Thompson, Barry C. (committee member)

Creator Email jannise.buckley@gmail.com,jjbuckle@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c40-209602

Unique identifier UC11276458

Identifier etd-BuckleyJan-4107.pdf (filename),usctheses-c40-209602 (legacy record id)

Legacy Identifier etd-BuckleyJan-4107.pdf

Dmrecord 209602

Document Type Dissertation

Format application/pdf (imt)

Rights Buckley, Jannise J.

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