University of Southern California Dissertations and Theses

Solution‐phase synthesis and deposition of earth‐abundant metal chalcogenide semiconductors

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Solution‐phase synthesis and deposition of earth‐abundant metal chalcogenide semiconductors

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SOLUTION-PHASE SYNTHESIS AND DEPOSITION OF EARTH-ABUNDANT
METAL CHALCOGENIDE SEMICONDUCTORS

by

Priscilla D. Antunez

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

August 2014

Copyright 2014 Priscilla D. Antunez
ii

Table of Contents

Acknowledgements ................................................................................................... iv

List of Figures ............................................................................................................. v

Abstract .................................................................................................................... viii

Chapter 1: Tin and Germanium Monochalcogenide IV-VI Semiconductor
Nanocrystals for Use in Solar Cells .......................................................................... 1
1.1 Abstract ...................................................................................................... 1
1.2 Introduction ................................................................................................ 1
1.3 General Structural Characteristics of Tin and Germanium
Monochalcogenides .................................................................................... 5
1.4 Nanocrystal Synthesis and Characterization .............................................. 9
1.4.1 Tin Sulfide Nanocrystals ............................................................. 9
1.4.2 Tin Selenide Nanocrystals ........................................................ 17
1.4.3 Tin Telluride Nanocrystals ....................................................... 25
1.4.4 Germanium Sulfide and Selenide Nanocrystals ....................... 30
1.5 Photovoltaic Device Applications ............................................................ 31
1.6 Conclusions .............................................................................................. 35
1.7 References ................................................................................................ 37

Chapter 2: Low Temperature Solution-Phase Deposition of SnS Thin Films ... 48
2.1 Introduction .............................................................................................. 48
2.2 Results and Discussion ............................................................................ 50
2.2.1 Dissolution and Deposition Methodology ................................ 50
2.2.2 Structural and Optical Characterization .................................... 52
2.2.3 Photocurrent Response Measurements ..................................... 54
2.3 Conclusion and Future Work ................................................................... 57
2.4 Experimental Details ................................................................................ 57
2.4.1 General Considerations ............................................................. 57
2.4.2 Dissolution of bulk SnS ............................................................ 58
2.4.3 SnS Recovery and Organic Content Determination ................. 58
2.4.4 Structural and Optical Characterization .................................... 58
2.4.5 Spin-Coating ............................................................................. 60
2.4.6 Photoelectrochemical Characterization .................................... 61
2.5 Acknowledgements .................................................................................. 61
2.6 Refences ................................................................................................... 61

iii

Chapter 3: Solution-Phase Synthesis of Highly Conductive Tungsten Diselenide
Nanosheets ................................................................................................................ 64
3.1 Introduction .............................................................................................. 64
3.2 Results and Discussion ............................................................................ 66
3.2.1 Synthetic Methodology ............................................................. 66
3.2.2 Structural Characterization ....................................................... 68
3.2.3 Electrical Transport Properties ................................................. 74
3.3 Conclusions and Future Work ................................................................. 75
3.4 Experimental Details ................................................................................ 75
3.4.1 General Considerations ............................................................. 75
3.4.2 Synthesis of WSe
2
nanosheets .................................................. 76
3.4.3 Structural Characterization ....................................................... 77
3.4.4 Device Characterization ............................................................ 78
3.5 Acknowledgements .................................................................................. 79
3.6 References ................................................................................................ 79

Chapter 4: Low Temperature Solution-Phase Deposition of SnSe Nanocrystalline
Thin Films ................................................................................................................. 84
4.1 Introduction .............................................................................................. 84
4.2 Results and Discussion ............................................................................ 86
4.2.1 Synthetic Methodology and Nanoparticle Characterization ..... 86
4.2.2 Deposition and Ligand Exchange Methodology ....................... 89
4.2.3 Optical Thin Film Characterization .......................................... 91
4.2.4 Photocurrent Response Measurements ..................................... 93
4.3 Conclusion and Future Work ................................................................... 95
4.4 Acknowledgements .................................................................................. 96
4.5 References ................................................................................................ 96

Chapter 5: Solution-Phase Synthesis of NiS Nanocrystals ................................... 99
5.1 Introduction .............................................................................................. 99
5.2 Experimental Details .............................................................................. 101
5.2.1 General Considerations ........................................................... 101
5.2.2 Nanocrystal Syntheses ............................................................ 101
5.2.3 Nanocrystal Characterization .................................................. 102
5.3 Results and Discussion .......................................................................... 102
5.3.1 NiS Micro- and Nanocrystals .................................................. 102
5.3.1.1 Structural Characterization ...................................... 103
5.3.2 Ni
9
S
8
Nanocrystals .................................................................. 105
5.3.2.1 Structural Characterization ...................................... 106
5.4 Conclusion ............................................................................................. 107
5.5 Acknowledgements ................................................................................ 108
5.6 References .............................................................................................. 108

Bibliography ........................................................................................................... 111
iv

Acknowledgements
I would like to thank my research advisor Prof. Richard Brutchey, as well as professors
Mark Thompson, Hannah Reisler, and Suketu Bhavsar for their continued support and
mentorship. I also thank the members of the Brutchey and Thompson laboratory groups
for their fellowship and support, as well as the rest of the USC community. I thank my
grandmother Yolanda Abelmann for her prayers; Rachel and Anselmo Garza for their
support; and my parents Yolanda Roman de Antunez and Octavio Antunez for their
unconditional encouragement. I would like to especially thank my loving husband
Andres Garza for his immense patience and support.
I would also like to thank Michele Dea, Susan Peterson, Phillip Sliwoski, Dr. Elizabeth
Erickson, and Elizabeth Bakus for all their help. I thank Jose Araujo, Dr. David Webber,
Dr. Elsa Couderc, and Jannise Buckley for the many hours of hard work, help, and
friendship. I also acknowledge the invaluable help and training from Ernesto Barron,
John Curulli, Frank Devlin, and Doug Houser.
Lastly, I would like to thank my undergraduate advisors, Dr. Suketu Bhavsar, Dr. Maria
Botero-Omary, Dr. Laurie Starkey, Dr. Barbara Burke, and Dr. John Edlund for their
continued support and mentorship.
v

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 ....................................... 12

Figure 1.5 Photocurrent response of a thin film of SnS nanocrystals ................. 13
Figure 1.6 SEM image of 6-nm SnS nanocrystals .............................................. 15

Figure 1.7 TEM images of SnSe nanocrystals .................................................... 20

Figure 1.8 Low- and high-resolution TEM image 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 ......................................... 26

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 ............ 33

Figure 2.1 Dissolved bulk SnS in ethylenediamine and 1,2-ethanedithiol ......... 50

Figure 2.2 Thermal gravimetric analysis for a SnS solution ............................... 51

Figure 2.3 FT-IR spectra and powder XRD of SnS ink ...................................... 52

Figure 2.4 Rietveld analysis of the XRD pattern of SnS nanocrystals ................ 53

Figure 2.5 Diffuse reflectance UV-vis-NIR spectroscopy .................................. 54

Figure 2.6 SEM images of SnS films .................................................................. 55
Figure 2.7 Photocurrent response of a typical SnS thin-film electrode ............... 56
vi

Figure 3.1 Hexagonal molybdenite structure ...................................................... 65

Figure 3.2 The WSe
2
nanosheets regular and SEM image .................................. 67
Figure 3.3 TGA analysis of WSe
2
nanosheets .................................................... 67

Figure 3.4 XRD patterns of WSe
2
nanosheets .................................................... 69

Figure 3.5 Raman spectra of a WSe
2
................................................................... 69

Figure 3.6 SEM-EDX elemental analysis of a WSe
2
nanosheet sample ............. 70

Figure 3.7 HR-TEM images of typical WSe
2
nanosheets ................................... 71
Figure 3.8 HR-TEM of unannealed TEM samples dried from toluene ............... 71

Figure 3.9 Typical TEM image of WSe
2
nanosheets dried from TMU .............. 72
Figure 3.10 Low resolution TEM of isolated, unannealed WSe
2
nanosheets ....... 72

Figure 3.11 SAED and HR-TEM of unannealed WSe
2
nanosheets ...................... 73
Figure 3.12 Room temperature two-point I-V measurements ............................... 74

Figure 4.1 SnSe nanocrystals dispersed in toluene and TMU ............................. 87
Figure 4.2 Thermal gravimetric analysis ............................................................. 87

Figure 4.3 TEM images of a typical batch of SnSe nanocrystals ........................ 88
Figure 4.4 Powder XRD pattern of SnSe Nanocrystals ...................................... 89
Figure 4.5 FT-IR staggered spectra of a hydrazine treated SnSe thin film ......... 90

Figure 4.6 FT-IR staggered spectra of an ethanedithiol treated SnSe thin film .. 90

Figure 4.7 SEM images of the SnSe/TMU derived films ................................... 91

Figure 4.8 Absorption spectrum of a hydrazine treated SnSe thin film .............. 92

Figure 4.9 Absorption spectrum of an ethanedithiol treated SnSe thin film ....... 93

vii

Figure 4.10 SnSe nanocrystals spin coated on ITO with a constant applied potential,
472 nm illumination, before and after hydrazine treatment ............... 94

Figure 4.11 SnSe Nanocrystals spin coated on ITO with a constant applied potential,
472 nm illumination, before and after edt treatment. ......................... 95

Figure 5.1 XRD pattern of β-NiS (millerite) ..................................................... 103
Figure 5.2 XRD pattern of β-NiS (millerite) ..................................................... 104
Figure 5.3 TEM images of NiS dried from toluene .......................................... 105
Figure 5.4 XRD diffraction pattern of Ni
9
S
8
crystals ........................................ 106
Figure 5.5 TEM images of Ni
9
S
8
crystals and nanorods. .................................. 107

viii

Abstract
The solution phase deposition of inorganic semiconductors is crucial for the commercial-
scale fabrication of thin-film based devices such as photovoltaics, thermoelectrics, and
field effect transistors. The need for a truly dissolved inorganic semiconductor is
highlighted by inexpensive high throughput processes such as spray coating and roll-to-
roll printing that allow for large area and flexible substrate deposition, which are
especially desirable for photovoltaic device commercialization. We have addressed the
challenge of creating inks of several inorganic semiconductors via two low temperature
techniques: 1) the dissolution and deposition of bulk tin sulfide (SnS) thin films using a
hydrazine-free solvent mixture and 2) the solution-phase synthesis of tungsten selenide
(WSe
2
), tin selenide (SnSe), and nickel sulfide (NiS). The indirect band gaps of SnS,
WSe
2
, and SnSe (E
g
= 1.1 eV, E
g
= 1.2 eV, E
g
= 0.9 eV, respectively) make them
attractive for light harvesting applications. In addition, SnS has been reported as a highly
efficient photoelectrode material for water splitting, while WSe
2
has been incorporated
into highly efficient photoelectrochemical devices (PCE = 17%). The methods we have
developed allow for the facile solution-phase deposition of films by spin-coating for SnS
and SnSe or by evaporation for WSe
2
and NiS. The materials are shown to be phase pure
via powder XRD, SEM-EDS, diffuse reflectance UV-vis-NIR, Raman, and high
resolution TEM; while the organic content in the films is shown to be greatly reduced
using FT-IR. The SnS and SnSe thin film electrodes exhibit a p-type electrochemical
photoresponse, and the films made with WSe
2
nanosheets show a high conductivity of up
to 92 S cm
-1
after annealing.
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 processible.
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.
363 terawatt-hours per day in just seconds, but the high cost of current silicon-based
2

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 nanocrystal size and shape through the manipulation of
synthetic variables, such as reaction temperature, reaction time, reaction solvent, and
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 devices because
of the high level of synthetic control that has been achieved in these systems.
27-29

3

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/aluminum 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-
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
4

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 is 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).
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 bodyweight per day.
43
The
decreasing threshold for BLLs observed over the past 20-30 years underscores the
importance of exploring alternative semiconductor materials that are less toxic, but have
otherwise similar properties to the lead chalcogenides. Tin and germanium are less toxic
than lead; and unlike lead, inorganic tin compounds are not readily absorbed into the
blood stream through ingestion or inhalation.
44
Although studies have cautioned over
5

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 possess 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.
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.
6

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 (Figure 1.1). 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 (Figure 1.2).
51,52
The strong anisotropic properties
observed in SnS and SnSe, for example, have been correlated to this double layer
structure, and it has 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
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 =
7

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,
55a
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.
55b

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 Cmcm β-SnSe structure (a = 4.31 Å, b = 11.70 Å, and c =
4.31 Å).
54,56
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
congruent melt occurs.
57
Solid-state studies have shown a low temperature cubic β-SnTe
to rhombohedral α-SnTe phase transition occurs at ca. -200 ˚C.
54

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

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
8

600 ˚C, the orthorhombic Pnma structure undergoes a phase transition to the hexagonal
β-GeS structure (a = 8.70 Å, c = 8.73 Å).
58
Similarly, the thermodynamically preferred
phase of GeSe at room temperature is also an orthorhombic Pnma α-GeSe structure (a =
4.40 Å, b = 3.85 Å, and c = 10.82 Å).
59
This phase undergoes a phase transition at ca.
586 ˚C to the cubic Fm3m β-GeSe structure (a = 5.73 Å).
60

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)
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+
.
49,61-64
The Sn(5s) lone pair
9

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 (E
g
= 1.6–1.7 eV for GeS and E
g
= 1.1–1.2 eV for GeSe) that overlap well with the
solar spectrum.
65,66
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.
67,68
As a
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.
69,70
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.
71-76
In 2006, Greyson et al. synthesized SnS nano- and
10

microcrystals with a purported metastable zinc blende crystal structure, which has been
shown to be a pseudotetragonal phase.
76,55b
The SnS nano- and microcrystals were
prepared by the reaction of SnCl
2
with elemental sulfur at 170 ˚C in oleylamine.

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), which has been described instead as a
pseudotetragonal structure. Reprinted with permission from ref. 76 (Copyright 2006 Wiley-VCH).

Scanning electron microscopy (SEM) images showed tetrahedral nano- and microcrystals
ranging between 200–300 nm on each side (Figure 1.3a). Modifications in the reaction
time or amine surfactant (decyl-, dodecyl-, hexadecyl-, and oleylamine) did not have an
11

effect on the size, polydispersity, or shape of the SnS product. The powder X-ray
diffraction (XRD) was used to identify a metastable zinc blende crystal structure (a =
5.85 Å) with small amounts of Sn, sulfur, Sn(OH)
2
, and orthorhombic SnS (Figure 1.3b),
but Schaak et al. have instead described it as a crystallographic distortion of the Pnma
space group into a single pseudotetragonal structure (a = 11.55 Å, b = 4.12 Å, and c =
4.12 Å).
55b
The zinc blende phase had also been reported in bulk SnS grown epitaxially
on a NaCl seed layer,
77,78
and it had been thought to be mixed with the orthorhombic α-
SnS phase; the Schaak group explained how these patterns are all best described as a
single pseudotetragonal phase with some SnO
2
impurities.
55b
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). 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
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%) 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
12

zinc blende SnS predicts that the material is either a metal or a small indirect band gap
semiconductor; however, the absorption profile of the purported zinc blende tetrahedra is
blue shifted to 700 nm (E
g
= 1.8 eV) relative to the orthorhombic phase. The group
hypothesized that the higher energy band gap relative to the orthorhombic phase was a
result of the different symmetries of the crystal structures. The impurities observed by
XRD in the zinc blende SnS were also thought to play a role in the spectral blue shift of
the material.

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. 79 (Copyright 2008 American Chemical Society).

In 2008, Hickey et al. synthesized monodisperse and slightly sulfur-rich sub-10 nm SnS
nanocrystals with a narrow 10% size distribution (Figure 1.4).
79
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 1.5. 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.
79 (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 other phases 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
14

are relatively close in energy (i.e. 200–600 meV), the authors suggest that the positions 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 (Figure
1.5). It was found that the photocurrent response for the SnS nanocrystals was stable and
repeatable over many light/dark cycles.
In 2009, Xu et al. synthesized sub-5 nm SnS nanocrystals that were dispersible in polar
solvents, such as ethanol.
80
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
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.
15

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.

Figure 1.6. SEM image of 6-nm SnS nanocrystals synthesized from SnCl
2
and S(SiMe
3
)
2
.
Reprinted with permission from ref. 81 (Copyright 2010 IOP Publishing).

In 2010, Liu et al. synthesized SnS nanocrystals that were single crystalline,
monodisperse, and size-tunable (Figure 1.6).
81
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.
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.;
79

however, Liu et al. assigned it to a direct band gap transition rather than and indirect band
gap transition.
More recently, Ning et al. reported the synthesis of SnS nanostructures with various sizes
and morphologies.
82
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-
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.
17

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 340–380 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.
84,85
The indirect and direct band gaps (E
g
= 0.90 eV and 1.30 eV,
respectively) of bulk SnSe correlate well with the optimum band gap values for solar
cells, which fall between E
g
= 1.0–1.5 eV.
49
In 2007, Pejova et al. 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 an indirect and direct bandgap of E
g
= 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
18

1.65 eV, respectively. Promptly after demonstrating quantum confinement effects in
nanocrystalline SnSe thin films, Pejova et al. went on to demonstrate the potential of
SnSe as an active component in PVs by studying the charge-transport properties of these
films.
87
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 can act as an acceptor 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.
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.
88-91
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.
88
In 2000, Zhang et al. published an aqueous route to nanocrystalline SnSe with
large, sheet-like morphologies.
89
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)
19

and a narrow size distribution.
90
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.
89,91

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.
In 2002, Schlecht et al. took a different approach to the solution-phase SnSe synthesis by
employing diorganodichalcogenides as soluble chalcogenide sources.
92
They turned to
the 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 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
selenoate rather than the intended
20

SnSe. In the second step, thermolysis of Sn(SePh)
4
at 300 ˚C lead 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.

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. 93 (Copyright 2010 American Chemical Society).

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.
93
Our synthetic route
involved the use of a diorganodichalcogenide as the chalcogen source.
94-99
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
21

reaction temperature was raised to 180 ˚C for 4 min and then quenched by cooling to
obtain phase-pure SnSe nanocrystals. We 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.
92

Transmission electron microscopy analysis revealed the SnSe product to be composed of
elongated anisotropic nanocrystals of variable length and consistent width (19.0 ± 5.1
nm; Figure 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
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) that 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.
100
They successfully carried out a solution-
phase synthesis of SnSe through hot injection (65-175 ˚C) of Sn[N(SiMe
3
)
2
]
2
into TOPSe
in the presence of oleylamine. After nucleation, oleic acid was introduced to the mixture
22

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.
20,101
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
size was reached. Replacing oleic acid with dodecanethiol resulted in the growth of
anisotropic SnSe nanocrystals, similar to those produced by Franzman et al.
93
With this synthetic method, SnSe nanocrystals from 4 to 10 nm could be controllably
synthesized by tuning the injection and reaction temperatures (Figure 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
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.
23

This is an interesting result because a metastable phase is being observed at lower
temperatures.
54,56,102,103a

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. 100 (Copyright 2010 American Chemical Society).

Nanocrystals in the size range from 4–9 nm exhibited quantum confinement affects,
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 (Figure 1.9a), respectively, while
the direct band gap varied from E
g
= 1.8–1.3 eV over the same size range. Both the
indirect and direct band gaps demonstrated a rough 1/r
2
dependence on nanocrystal size
(Figure 1.9b). Although a prototype solar cell was not reported, the Hanrath group did
report the applicability of their SnSe nanocrystals as a photoconductive material. The
24

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.

Figure 1.9. (a) Tauc plot from UV-vis-NIR data for SnSe nanocrystals showing an indirect band
gap. (b) Approximate 1/r
2
relationship between band gap and mean nanocrystal size. Reprinted
with permission from ref. 100 (Copyright 2010 American Chemical Society).

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.
25

demonstrated control over shape, lateral dimensions, and sheet thickness by reacting
SnCl
2
, trioctylphosphine selenide (TOP-Se), and hexamethyldisilazane (HMDS) in
oleylamine using a one-pot approach.
103b
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 E
g
= 1 eV and preliminary photoresponse measurements were
determined from SnSe nanosheet samples that were drop casted.
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.
104,105
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.
106
Tin telluride is a direct band gap semiconductor that exhibits a narrow band gap of
E
g
= 0.2 eV at room temperature,
96
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.
108
One of the first solvothermal syntheses of tin monochalcogenides entailed the use of
soluble diaryl dichalcogenides and activated Sn
0
nanocrystals.
92
This work evolved from
the use of elemental selenium and tellurium to the use of diphenyl diselenide (Ph
2
Se
2
)
26

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.
109
The reaction of Sn
0
with Ph
2
Te
2

produced two different morphologies of 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
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. 108 (Copyright 2007 American Chemical
Society).

In 2007, Kovalenko et al. synthesized monodisperse SnTe nanocrystals in solution that
were tunable in size between 4.5–15 nm (Figure 1.10), with their corresponding band
27

gaps ranging between E
g
= 0.8–0.4 eV, respectively.
108
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
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.
110
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 (Figure 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.
28

More recently, Ning et al. were able to synthesize SnTe nanocrystals and nanowires.
111

The unique properties exhibited by nanorods and nanowires make these low dimensional
materials attractive,
112
and have motivated studies on their formation via aggregation of
0-D nanocrystals by oriented attachment.
113
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
crystallinity; however, these nanocrystals tranformed 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.
114a
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.
The application or further investigations of tin chalcogenide nanocrystalline structures
should carefully consider Reiss et al.’s findings on surface oxidation, as charge transfer
and transport are two processes that are strongly influenced by the surface quality of
nanocrystals. The group’s
119
Sn-Mössbauer spectroscopy results highlight the need for
the careful processing in an inert atmosphere and further surface engineering efforts of tin
29

chalcogenide nanocrystals.
114b
The group showed 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.

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

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.
115,116
Vaughn et al. have recently reported the first solution chemistry route for
colloidal GeS and GeSe nanostructures.
117
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 (E
g
= 1.6–1.7 eV for GeS and E
g
= 1.1–1.2 eV
for GeSe) that overlap well with the solar spectrum.
65,66
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
hexagons (2–4 µm by 0.5–1 µm; Figure 1.12), with thicknesses between 3–20 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.
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,
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 x 10
-6
S cm
-1
for drop-cast
thin films of the GeSe nanosheets, while a two-point I-V measurements demonstrate the
31

p-type character of this material (Figure 1.13). The conductivity value is comparable to
other colloidal nanocrystalline semiconductor thin films.
118

Figure 1.12. (a) TEM images of GeS and GeSe (inset) nanosheets synthesized from GeI
4
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. 117 (Copyright 2010 American Chemical Society).

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 has
32

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 appears 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.
119
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
,
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.
120
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-phenylene
33

vinylene] (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 x 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.

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. 117
(Copyright 2010 American Chemical Society).

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.
121
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.
The first demonstration of PV application of a cell based on TiO
2
/SnS films was reported
in 2010 by Wang et al.
122
An electrochemical solar cell structure of FTO/Pt + electrolyte
+ SnS/TiO
2
/FTO yielded a high V
OC
= 0.471 V, a J
SC
= 0.30 mA cm
-2
, a power
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.
123
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.
93
Their device consisted of a SnSe:poly[2-
35

methoxy-5-(3’,7’-dimethyloctyoxy)-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 an 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 wt/wt, 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 likely a
results 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
36

into a hot solution of 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
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. Moreoever, 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 structure
69
), 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
37

removing the insulating ligands protecting the nanocrystals and replacing them with
small molecules, such as hydrazine or metal chalcogenide clusters.
124,125
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.
118
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.
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48

Chapter 2. Low Temperature Solution-Phase Deposition of
SnS Thin Films*
* Manuscript in preparation

2.1. Introduction
The solution phase deposition of inorganic semiconductors is crucial for the commercial-
scale fabrication of thin film based devices such as photovoltaics, thermoelectrics, and
field effect transistors.
1
The need for a truly dissolved inorganic semiconductor is
highlighted by inexpensive, high throughput processes such as spray coating and roll-to-
roll printing that allow for large area and flexible substrate deposition, which are
especially desirable for thin film photovoltaic device commercialization.
2
The challenge
of dissolving inorganic semiconductors has been addressed by using colloidal
nanocrystals
3-9
or relatively harsh solvents to dissolve the bulk.
10
A true inorganic
semiconductor ink is often not attainable via the nanocrystal route due to low colloidal
stabilities, but the use of hydrazine has proven to be a more straightforward way to
dissolve many inorganic semiconductors.
1

Mitzi et al. have used hydrazine in combination with excess chalcogen to dissolve
numerous metal chalcogenide semiconductors, obtaining thin films after mild annealing
(270–540 °C).
1
In 2004, his group used dissolved SnS
2
and SnSe
2
in freshly distilled
hydrazine to produce thin films (ca. 5 nm) that showed n-type transport, large current
densities (>10
5
A cm
-2
) and high mobilities (>10 cm
2
V
-1
s
-1
).
2
The thin films were spin
49

coated in air onto oxidized silicon substrates with mild final annealing treatments at 270–
296 °C.
2
Furthermore, they have applied this approach to dissolve CuS
2
, In
2
Se
3
, Ga, and
Se to produce highly efficient CIGS solar cells (PEC = 15.2%, FF = 0.75, V
OC
= 623 mV,
J
SC
= 32.6 mA cm
-2
), which were annealed at 540 °C.
11
However, the recent pressure on
industry to adopt inherently safer technologies requiring the substitution of hazardous
chemicals with safer ones
12
highlights the importance of new hydrazine-free solvent
mixtures that can effectively dissolve inorganic semiconductors without the use of highly
toxic, explosive, or carcinogenic solvents. Most recently, hydrazine-free solutions
consisting of ethylenediamine/alkanethiol mixtures have been used to dissolve and
synthesize several bulk inorganic semiconductors after driving off the solvent mixture via
mild annealing.
13,14
Bulk tin sulfide (SnS) is a semiconductor comprised of earth abundant elements, which
has a direct (E
g dir
= 1.32 eV) and indirect (E
g ind
= 1.08 eV) band gap that make it
attractive for light harvesting applications.
15
Tin sulfide is relatively stable and nontoxic,
it exhibits high carrier concentrations (1–3 × 10
18
cm
-3
),
15
high anisotropic hole mobilities
(90 cm
2
V
-1
s
-1
,
┴
to c axis),
15
and a high absorption coefficient (>10
4
cm
-1
).
16
The
material exhibits an orthorhombic (Pnma) herzerbergite structure. Tin sulfide is an
intrinsic p-type semiconductor that can be doped to exhibit n-type behavior upon
annealing in air.
17,18
In addition, SnS can be doped with Ag, Al, Ni, and Cl due in part to
its layered structure,
17
which is composed of double layers that are weakly held together
via van der Waals forces perpendicular to the c axis. These double layers are comprised
of strongly bound double layers of Sn and S that have no dangling bonds.
15
The wide
50

variety of Sn–S phases (e.g. SnS, SnS
2
, Sn
2
S
3
, Sn
3
S
4
, Sn
4
S
5
)
15
can make it challenging to
obtain a phase pure material from a solution-phase route. Herein, we report the facile
fabrication of phase-pure SnS thin film photoelectrodes from the dissolution of bulk SnS
in ethylenediamine (en) and 1,2-ethanedithiol (edt), followed by recovery and
crystallization via a mild annealing step (350 °C) to drive off the organic solvent mixture.
2.2. Results and Discussion
2.2.1. Dissolution and Deposition Methodology
Ethylenediamine, along with alkanethiols, have proven to be useful solvent mixtures to
easily dissolve inorganic semiconductors.
13,14
Here, bulk SnS was easily dissolved in an
11:1 vol/vol chelating mixture of ethylenediamine (en) and 1,2-ethanedithiol (edt) under
nitrogen; thin films of SnS were in turn prepared by spin coating to demonstrate the
utility of the system for solution processing (Figure 2.1, right).

Figure 2.1. Dissolved bulk SnS in ethylenediamine and 1,2-ethanedithiol (edt) solvent mixture
(left) and a SnS thin film on a glass substrate (right).
51

The concentration of the dissolved species increased with prolonged stirring and heating
(120 mg mL
-1
maximum solubility), but most of the characterization was performed on
material stirred for ca. 15 h at 50 °C, which produced a 60 mg mL
-1
clear, light yellow
solution (Figure 2.1, left). Thermal gravimetric analysis (TGA) was used to determine
the temperature needed to recover the SnS material from solution. After drying the
solution in situ at 125 °C for 15 min, the sample was further heated to 425 °C. A mass
loss of ca. 50 wt% at 350 °C was assigned to the loss of organics (Figure 2.2), which
corresponded to an approximate concentration of 60 mg mL
-1
, or 6 wt% of the total
volume.

Figure 2.2. Thermal gravimetric analysis showing the decomposition endpoint for a SnS solution.

FT-IR spectroscopy corroborated the loss of organics after mild annealing of the SnS
solution to 350 °C under flowing nitrogen. Figure 2.3 (a) shows the strong ν(C–H) and
52

ν(N–H) stretching bands (ca. 3350–2750 cm
-1
), which disappear after annealing the dried
SnS solution to 350 °C. The absence of a ν(S–H) stretching band (ca. 2500–2600 cm
-1
)
is in agreement with our previous findings, which indicate the almost full deprotonation
of edt in the solvent mixture.
13
2.2.2. Structural and Optical Characterization
Powder X-ray diffraction (XRD) confirmed the recovered material was phase-pure SnS.
Oxide impurities were avoided by drying and annealing the SnS ink under flowing
nitrogen, while a mixture of SnS, SnS
2
, and SnO resulted when annealed in air (380 °C).
Figure 2.3 (b) shows the matching diffraction patterns of both the standard (commercial)
SnS material and that of recovered SnS after annealing (350 °C).

Figure 2.3. (a) FT-IR spectra of a SnS ink dried to 130 °C for 15 min and the same sample but
after annealing to 350 °C. (b) Powder XRD patterns of a commercial SnS sample (standard) and
that obtained after annealing the SnS ink to 350 °C under flowing nitrogen.

53

The diffraction patterns display a 100% intensity peak (2θ = 31.5°), indexed to the (111)
reflection of orthorhombic SnS (JCPDS no. 01-075-1803), while all other peaks in Figure
2.3 (b) also match the orthorhombic SnS phase (e.g. peaks at 22.2°, 27.6°, 39.2°, and
48.7° can be indexed to the (110), (021), (131), and (230) reflections, respectively).
Rietveld analysis was used to calculate the lattice constants for the recovered SnS sample
(a = 11.242(3) Å, b = 3.9943(11) Å, c = 4.3102(12) Å) that closely match literature
values of a = 11.20 Å, b = 3.99 Å, c = 4.33 Å (Figure 2.4).
19
In addition, scanning
electron microscope energy-dispersive X-ray spectroscopy (SEM-EDX) corroborated the
composition of the recovered material. Analysis of randomly selected areas of powdered
samples gave an average elemental composition of 53 at% Sn and 47 at% S for the SnS
reference, and 52 at% Sn and 48 at% S for the SnS recovered from solution, which are
within instrumental error.

Figure 2.4. Rietveld analysis of the XRD pattern of SnS nanocrystals. Experimental (×) and
calculated (⎯) patterns are shown along with the difference curve (⎯). Tickmarks (⏐)
corresponding to the phase refined are also shown.
54

The band gap of the recovered SnS material was determined via diffuse reflectance UV-
vis-NIR spectroscopy, which allowed for both the investigation of the material’s optical
properties and to further probe the phase purity of the material. Figure 2.5 shows the
typical wide absorption spectrum of recovered SnS (ca. 3%) mixed in BaSO
4
. The
Kubelka-Munk function was applied to calculate both the direct (E
g dir
= 1.3 eV) and
indirect (E
g ind
= 1.1 eV) band gaps (Figure 2.5 inset), which closely match the literature
(vide supra) and our reference experimental values of ca. E
g dir
= 1.3 ev and E
g ind
= 1.1
eV).

Figure 2.5. Diffuse reflectance UV-vis-NIR spectroscopy. Absorption spectrum from recovered
SnS and its corresponding indirect band gap (inset).
2.2.3. Photocurrent Response Measurements
The photocurrent response of the SnS thin films was measured to assess a material’s
potential for photovoltaic applications using a facile and nondestructive electrochemical
method.
20-24
Figure 2.7 shows the photocurrent results of a typical SnS electrode made
using two coats of a 60 mg mL
-1
SnS ink that was spin-coated (1250 rpm, 3 min) onto
55

FTO covered glass under flowing nitrogen. The films were then annealed to 350 °C,
which is the end-point of solvent decomposition determined by TGA and confirmed by
FT-IR (vide supra). A final 500 °C annealing step in a tube furnace under flowing
nitrogen increased the mechanical robustness of the thin film electrodes to conduct the
photoelectrochemical characterization.

Figure 2.6. SEM images of SnS films using two coats of a 60 mg mL
-1
SnS ink annealed to 350
°C after each layer, with one final anneal to 500 °C. A top-down view of the film shows no
major cracks or pin-holes over a large area (top, left), and its corresponding cross-section image
which suggests reduced interparticle spacing (top, right), when compared to a film annealed
quickly on a pre-heated hot plate (bottom).
56

SEM images show that the resulting films have no major cracks or pin holes even after
the final 500 °C anneal (Figure 2.6 left); the images also suggest that there is less
interparticle spacing with a lower heating rate of the interlayer annealing step when going
from room temperature to 350 °C (Figure 2.6 right), as opposed to the results when using
a preheated plate (Figure 2.6 bottom).

Figure 2.7. Photocurrent response of a typical SnS thin-film electrode exhibiting p-type behavior
under a bias of -700 mV (vs SCE).
The resulting electrodes exhibited p-type behavior in a three-electrode
photoelectrochemical cell with a Pt counter electrode and a SCE reference electrode in an
aqueous solution containing 0.1 M Eu(NO
3
)
3
as the redox mediator. Figure 2.7 shows the
current response under 1 sun chopped illumination (ELH lamp, 300 W). An increase in
57

current response of ca. 36 µA cm
-2
was observed at -500 mV, while a bias of -700 mV
produced an increase of almost 170 µA cm
-2
.
2.3. Conclusion and Future Work
In summary, we report on the facile deposition of phase-pure SnS thin films via the
dissolution of bulk SnS. The band gaps of the recovered material (E
g dir
= 1.3 eV and
E
g ind
= 1.1 eV) match that of the reference sample and the literature values.
Morphological film control was achieved by regulating the heating rate in the post-
deposition annealing step between layers, resulting in films devoid of major cracks and
pin-holes. Photocurrent response measurements showed that the films exhibit p-type
behavior with a strong and stable photocurrent response upon 1 sun illumination, which
highlights the potential of this material for photovoltaic devices. Future work will
examine the viability of this material for use in solid-state devices.
2.4. Experimental Details
2.4.1. General Considerations
Tin(II) sulfide (Aldrich, 99.99+%) and europium(III) nitrate hexahydrate (Strem,
99.9+%) were purchased and used without further purification. Both 1,2-
ethylenediamine (Fluka, 99.5+%) and 1,2-ethanedithiol (98+%, Alfa Aesar) were
distilled prior to use. 1,2-Ethylenediamine (en) was first dried for 3 d with either CaO
(~120 g per 500 mL) or CaH
2
(~15–20 g per 1 L); after decanting, en was refluxed over
Na until dark blue (ca. 6.5 h) to then be distilled under nitrogen. 1,2-Ethanedithiol (edt)
58

was dried over 3 Å molecular sieves, then distilled. The dissolution was performed using
standard Schlenk techniques under nitrogen in the absence of water and oxygen.
2.4.2. Dissolution of bulk SnS
The typical dissolution experiment entailed adding ca. 650 mg of thoroughly ground (15
min) SnS into a three-neck round-bottom flask fitted with a reflux condenser and a stir
bar under flowing nitrogen. Distilled en (10 mL) and edt (0.9 mL) were added. The
mixture was lightly heated with a heat gun and sonicated for a few minutes (ca. 10 min).
A heating mantle was then used to further heat the mixture to 50 °C for 15 h while
stirring.
2.4.3. SnS Recovery and Organic Content Determination
Thermogravimetric analysis (TGA) was performed in a TA Instruments Q50 TGA. An
alumina pan was used to dry 50 µL of the solution mixture in situ. The program entailed
heating to 125 °C and holding this temperature for 15 min, to then cool to 30 °C and
equilibrate. The final analysis step entailed heating the sample to 425 °C at 10 °C min
-1
.
Figure 2.2 shows the data from this final analysis step only. FT-IR spectroscopy was
performed in a Bruker Vertex 80v using a ZnSe substrate to drop cast the SnS mixture
and anneal on a hot plate under flowing nitrogen.
2.4.4. Structural and Optical Characterization
The powder X-ray diffraction (XRD) patterns of typical samples were collected using a
Cu Kα radiation source (λ = 1.5406 Å) on a Rigaku Ultima IV diffractometer. The
59

diffraction patterns were recorded between 10˚ and 80˚ at room temperature. Rietveld
Analysis. Rietveld structural refinements were carried out using the GSAS software.
23,24

Table 2.1. Structural Parameters of SnS Nanocrystals Extracted From Rietveld Analysis
a
(Å)
11.242(3)

b
(Å)
3.9943(11)

c
(Å)
4.3102(12)

V
(Å
3
)
193.54(16)

x
Sn

0.12356(8)

y
Sn

0.25

z
Sn

0.11164(14)

x
S

0.1348(3)

y
S

0.75

z
S

0.5199(4)

U
Sn

(×100)
1.33(4)

U
S

(×100)
1.98

R
w

(%)
8.72

χ
2

1.47

The crystal structure of SnS nanocrystals was refined with the orthorhombic Pnma space
group. 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 pseudo-Voight function,
25
(4) lattice
60

constants (a, b, and c), (5) fractional atomic coordinates of the tin (x
Sn
, y
Sn
, z
Sn
) and
oxygen atoms (x
S
, y
S
, z
S
) constrained by the site symmetry, and (6) isotropic displacement
parameter of the tin atom (U
Sn
). The isotropic displacement parameter of the sulfur atom
(U
S
) was fixed at 1.5 × U
Sn
. The R
wp
and χ
2
indicators were employed to assess the
quality of the refined structural models.
26
Scanning electron microscope-energy
dispersive X-ray spectroscopy (SEM-EDX) was used for elemental analysis on a JEOL
JSM-7001F scanning electron microscope; SEM imaging was also performed using this
instrument. Diffuse reflectance UV-Vis-NIR spectroscopy was performed using the
reflectivity mode in a Perkin-Elmer Lambda 950 equipped with a 150 mm integrating
sphere. The samples were prepared by thoroughly grinding 12 mg of SnS with 400 mg
BaSO
4
and using a powder sample holder at the end of the integrating sphere.
2.4.5. Spin-Coating
A solution of ca. 60 mg mL
-1
was filtered using a 0.45 µm filter as 18 drops were
deposited on an FTO/glass substrate (1 × 1 in
2
) inside a spin coater under flowing
nitrogen. The thin films were spin coated using a Laurell Technologies Corporation
WS400Ez-6NPP-LITE single wafer spin processor at 1250 rpm (3 min) for each layer, at
an acceleration of 770 rpm s
-1
for the first coat and 990 rpm s
-1
for the second coat. The
films were annealed between layers (ca. 5 °C min
-1
) on a hot plate from room
temperature to 360 °C under flowing nitrogen. A final anneal of 500 °C in a tube
furnace, with flowing nitrogen, was performed to increase film robustness for the
photoelectrochemical experiments.
61

2.4.6. Photoelectrochemical Characterization
The photocurrent response measurements on the SnS thin films were performed using a
Gamry Reference 600 potentiostat. The glass photoelectrochemical cell contained a Pt
counter electrode, an SCE reference electrode, and a SnS thin film on FTO covered glass
as the working electrode. The electrolyte consisted of aqueous 0.1 M Eu(NO
3
)
3
as the
redox couple. The solution was purged with nitrogen via bubbling for 10 min prior to the
experiments and maintained flowing over the head space during the experiments. A 1
sun illumination was provided by an ELH lamp calibrated using a Thor Labs Si
photodiode. The total illuminated area of SnS was ca. 1 × 1 cm
2
, which was obtained by
cutting 2 pieces from the center of the 1 × 1 in
2
samples. Silver paint was used to attach a
Sn wire, and Loctite 9460 epoxy was then applied over the wire and around the edges of
the film to mask these areas from solution.
2.5. Acknowledgements
Acknowledgement is made to Dr. David Webber, Daniel Torelli, Carrie McCarthy, Jose
Araujo, Dr. Federico Rabuffetti, Ian McFarlane, Dr. Fan Yang, Prof. Nathan Lewis, and
Prof. Mark Thompson for both experimental assistance and helpful discussions.
2.6 References
1. Todorov T.; Mitzi, D. B. “Direct liquid coating of chalcopyrite light-absorbing
layers for photovoltaic devices.” Eur. J. Inorg. Chem. 2010, 17.
2. Mitzi D. B.; Kosbar L. L.; Murray C. E.; Copel M; Afzali A. “High-mobility
ultrathin semiconducting films prepared by spin coating.” Nature 2004, 428, 299.
62

3. Antunez P. D.; Buckley J. J.; Brutchey R. L. “Tin and germanium
monochalcogenide IV-VI semiconductor nanocrystals for use in solar cells.”
Nanoscale, 2011, 3, 2399.
4. Kwon, S.; Hyeon, T. “Colloidal chemical synthesis and formation kinetics of
uniformly sized nanocrystals of metals, oxides, and chalcogenides.” Acc. Chem. Res.
2008, 41, 1696.
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Alivisatos, A. P. “Shape control of CdSe nanocrystals.” Nature 2000, 404, 59.
6. Liu, S.; Guo, X.; Li, M.; Zhang, W.-H.; Liu, X.; Li, C. “Solution-phase synthesis and
characterization of single-crystalline SnSe nanowires.” Angew. Chem. Int. Ed. 2011,
50, 12050.
7. Hillhouse, H. W.; Beard, M. C. "Solar cells from colloidal nanocrystals:
Fundamentals, materials, devices, and economics." Curr. Opin. Colloid Interface
Sci. 2009, 14, 245.
8. Burda C.; Chen X. B.; Narayanan R.; El-Sayed M. A. “Chemistry and properties of
nanocrystals of different shapes.” Chem. Rev. 2005, 105, 1025.
9. Murray C. B.; Kagan C. R.; Bawendi M. G. “Synthesis and characterization of
monodisperse nanocrystals and close-packed nanocrystal assemblies.” Annu. Rev.
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10. Mitzi D. B.; Yuan M.; Liu W.; Kellock A. J.; Chey S. J.; Deline V.; Schrott A. G. “A
hig-efficiency solution-deposited thin-film photovoltaic device.” Adv. Mater. 2008,
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Cu(In,Ga)(S,Se)
2
absorber yielding a 15.2% efficient solar cell.” Prog. Photovolt.
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chalcogenides: Dissolution of
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ZnSnS
4
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64

Chapter 3. Solution-Phase Synthesis of Highly Conductive
Tungsten Diselenide Nanosheets*
*Published in part in Chemistry of Materials, 2013, 25, 2385–2387.

3.1. Introduction
Early transition metal dichalcogenides are an important class of layered materials that
have been used for hydrogen storage,

transistors, lubricants, catalysis, capacitors,
batteries, and photovoltaic devices.
1-8
The structure of transition metal dichalcogenides
consists of repeating crystalline layers, which are responsible for their 2D anisotropic
physical properties. Notable properties of transition metal dichalcogenides include
superconducting behavior in NbSe
2
and TaS
2
,
9-10
and the observation of ultralow thermal
conductivities (0.05 W m
–1
K
–1
) in disordered WSe
2
crystals.
11
In the nanoscale regime,
transition metal dichalcogenides have recently been shown to exhibit properties that
make them particularly interesting when compared to bulk material. For example, MoS
2

nanotubes are more air stable and exhibit greater loading capacities of Li than in the
bulk.
12
Moreover, certain transition metal dichalcogenides have been shown to change
from indirect to direct band gap semiconductors when isolated as single layers.
13-15
Bulk tungsten diselenide (WSe
2
) is an indirect semiconductor that can be doped to
generate a p- or n-type type material.
16-17
Bulk WSe
2
possesses a band gap of E
g indir
=
1.2 eV,
18
a high absorption coefficient, and high photostability.
16
In addition, its
incorporation into photoelectrochemical devices has resulted in conversion efficiencies
65

up to 17% (J
SC
= 38 mA cm
–2
).
19
Bulk WSe
2
exhibits hole carrier concentrations of
3.88 10
17
cm
–3
and highly anisotropic conductivities that range between 0.7 and 6.0 S
cm
–1
for in-plane measurements.
20-21
Tungsten diselenide crystallizes in Se-W-Se layers
where tungsten coordinates through strong covalent-ionic bonds (ca. 13% fractional ionic
character) to six selenium atoms in a trigonal prismatic geometry.
22
Variations in layer
stacking results in different polytypes; the 2H-WSe
2
polytype refers to the two layers and
the hexagonal structure represented in one unit cell (Figure 3.1), corresponding to the
hexagonal molybdenite structure (P6
3
/mmc).
23

Figure 3.1. Hexagonal molybdenite structure corresponding to the 2H-WSe
2
polymorph,
composed of Se-W-Se layers. The repeating crystalline layers are responsible for the material’s
2D anisotropic physical properties.

Traditional methods to prepare WSe
2
include sputtering, chemical vapor transport, solid-
state reactions, and electrodeposition;
24-28
however, there have been very few reports of
66

solution-phase syntheses of colloidal WSe
2
nanostructures. In 2000, Huang et al.
synthesized 4-7 nm WSe
2
nanoparticles by the reaction of H
2
Se with WCl
4
dissolved in a
ternary tridodecylmethylammonium iodide/hexane/octane inverse micelle solution.
29
In
2004, Duphil et al. synthesized WSe
2
nanoparticles by the reaction of elemental Se with
W(CO)
6
dissolved in p-xylene at 140 °C for several hours.
30
The resulting 10-30 nm
nanoparticles were then annealed at 550 °C, with XRD patterns revealing crystalline
tungsten oxide impurities. Herein, we present the first report of a facile solution-phase
synthesis of colloidal WSe
2
nanosheets.
3.2. Results and Discussion
3.2.1. Synthetic Methodology
Diorganodichalcogenides have proven to be useful chalcogen sources for the facile
solution-phase synthesis of colloidal semiconductor nanocrystals.
31-36
The synthesis of
WSe
2
nanosheets was achieved by the injection of 185 µL (0.92 mmol) di-tert-butyl
diselenide (
t
Bu
2
Se
2
) into a solution of 150 mg (0.46 mmol) WCl
4
in 25 mL dodecylamine
at 150 °C under nitrogen. The solution was then heated to 225 °C and held at this
temperature for 6 h prior to quenching. Tetra-n-octylammonium bromide (TOAB) was
added during the work-up to prevent agglomeration of the nanosheets, following an
adapted version of a previously published method.
37
The final washed product was
highly dispersible in tetramethylurea (TMU) and 1,2-dichlorobenzene (Figure 3.2, Left),
forming colloidal suspensions that were stable over the course of several months.
Although dispersions in toluene were not very colloidally stable, they produce visually
67

smooth films upon drop casting that were used throughout this study unless otherwise
noted. The omission of TOAB during the work-up procedure yielded macroscopic,
scroll-like pieces of WSe
2
instead of films upon deposition (Figure 3.2, Right).

Figure 3.2. The WSe
2
nanosheets were highly dispersible in 1,2-dichlorobenzene (Left). SEM
image of WSe
2
nanosheets without the use of TOAB in the work-up procedure shows
macroscopic cracking of the material (Right). Scale bar represents 500 µm.

Figure 3.3. TGA analysis of WSe
2
nanosheets shows mass loss above 200 °C and leveling off
below 500 °C (black curve). DSC shows two endothermic peaks that are likely due to loss of
organic material (green curve).
68

The synthesis was reproducible and yielded ca. 140 mg WSe
2
. Based on tungsten, the
WSe
2
yield was ~90%, taking into account the 12% organic content observed by
thermogravimetric analysis (TGA) (Figure 3.3). Differential scanning calorimetry (DSC)
corroborated the endothermic loss of organic material between 250-475 °C (Figure 3.3).
3.2.2. Structural Characterization
Powder X-ray diffraction (XRD) patterns of annealed and unannealed WSe
2
nanosheets
both appear to be phase pure without any crystalline tungsten oxide impurities in the 20-
30° 2θ range (Figure 3.4). Despite the air stability of the product, tungsten oxide
impurities were readily obtained unless strictly anaerobic conditions were employed
during both the synthesis and annealing steps. The three XRD patterns in Figure 3.4
display a 100% intensity peak (2θ = 13.5°) indexed to the (002) reflection of the 2H-
WSe
2
phase (JCPDS no. 00-038-1388).
38
Unannealed, drop-cast films displayed the
(002) reflection as the prominent diffraction peak along with very small peaks at ca. 2θ =
42° and 57° that correspond to the (006) and (008) reflections, respectively, suggesting a
strong preference for [001] orientation. Powdering the sample to reduce preferred
orientation effects resulted in a diffraction pattern showing very broad diffraction peaks
between 2θ = 30-60°, including the (102) reflection. Upon annealing the powdered WSe
2

nanosheets in a tube furnace for 5 min (475 °C under flowing nitrogen), additional peaks
became prominent; these were indexed to the (101), (103), and (105) lattice planes. The
Raman spectra of the annealed and unannealed WSe
2
films were collected to corroborate
the X-ray structure data.
69

Figure 3.4. XRD patterns of WSe
2
nanosheets before and after annealing at 475 °C.

Figure 3.5. Raman spectra of a WSe
2
powder from Alfa Aesar (blue), annealed (green) and
unannealed (gray) WSe
2
nanosheet films on glass; peaks highlighted by asterisks match both the
spectrum of the “as bought” material and that of reported single crystal WSe
2
.
40

70

The Raman spectra of annealed and unannealed WSe
2
films both reveal a characteristic
band at ca. 257 cm
–1
that can be assigned to the almost degenerate E
1
2g
and A
1g
Raman
active modes of WSe
2
, while the band at ca. 179 cm
–1
was assigned to the E
1g
mode,
indicative of the 2H-WSe
2
structure (Figure 3.5).
39
The unassigned peaks at ca. 127 and
319 cm
–1
have been observed in WSe
2
single crystals
40
and are also present in the Raman
spectrum of “as bought” WSe
2
powder (Alfa Aesar).

Figure 3.6. SEM-EDX elemental analysis of a WSe
2
nanosheet sample on a Si substrate. (a)
unannealed and (b) annealed to 475 °C. The composition for the unannealed sample was 10.9%
Se and 5.4% W, while the composition for the annealed sample was 18.7% Se and 10.4% W.

Scanning electron microscope energy-dispersive X-ray spectroscopy (SEM-EDX) and
ion coupled plasma atomic emission spectrometry (ICP-AES) were used to analyze the
elemental composition of the resulting WSe
2
nanosheets. Analysis of randomly selected
areas gave an average W/Se composition of 1:2.0 for the unannealed material and a 1:1.8
for the annealed material (Figure 3.6). ICP-AES results were in close agreement, giving
an average W/Se composition of 1:2.3 for the unannealed WSe
2
nanosheets.
71

Figure 3.7. HR-TEM images of (a) typical WSe
2
nanosheets dried from TMU, (b) nano-onion
structures of WSe
2
nanosheets dried from toluene, and (c) lattice fringes with interplanar spacings
of d = 2.85 and 2.72 Å. (d) SAED pattern of annealed WSe
2
nanosheets indexed to 2H-WSe
2
.

Figure 3.8. HR-TEM of unannealed TEM samples dried from toluene showing some
agglomerates that formed (a) nano-onion structures, and (b) long, aligned plates.
72

Figure 3.9. (a) Typical TEM image of WSe
2
nanosheets dried from TMU. (b) TEM image
showing the lattice fringes measured and averaged inside the white box. (c) The average spacing
of the fringes depicted within the box in (b) is d = 6.6 Å, corresponding to the (002) lattice plane.

Figure 3.10. Low resolution TEM of isolated, unannealed WSe
2
nanosheets. (a) Unannealed
WSe
2
nanosheets dispersed and dried from toluene with excess TOAB. (b) Unannealed WSe
2

nanosheets dispersed and dried from toluene.

73

Transmission electron microscope (TEM) analysis of the product revealed the nanosheet
morphology of the WSe
2
material. Samples drop-cast from toluene showed some
agglomerates with nano-onion structures and sheets with extended alignment (Figure
3.7b, Figure 3.8), while those dried from TMU showed less extensive alignment (Figure
3.7a, Figure 3.9). Individual thin sheets of various sizes were found around the
agglomerates (Figure 3.10). A high resolution TEM (HR-TEM) image of a nanosheet
displaying its (100) and (101) lattice planes (d = 2.85 and 2.72 Å, respectively) is shown
in Figure 3.7c.

Figure 3.11. (a) SAED of unannealed WSe
2
nanosheets. (b) HR-TEM of unannealed WSe
2

nanosheets dispersed and dried from TMU. (c) Lattice fringe analysis with an average spacing of
d = 2.8 Å, corresponding to the (101) lattice planes.

74

Although lattice fringes were observed throughout the unannealed WSe
2
nanosheets, only
a few diffuse rings were observed by selected area electron diffraction (Figure 3.11).
This suggests that the unannealed WSe
2
nanosheets are weakly crystalline, as the powder
XRD data also suggests. In contrast, more intense diffraction was observed in the SAED
pattern of the annealed WSe
2
nanosheets; these were indexed in agreement with the XRD
results (Figure 3.7d).
3.2.3. Electrical Transport Properties
The electrical transport properties of the solution-processed WSe
2
nanosheets were also
studied. Figure 3.12 shows the current-voltage (I-V) characteristics of an unannealed
WSe
2
film that was solution deposited between two aluminum electrodes. The material
was simply drop-cast from toluene and allowed to dry in air to give [001] oriented films
(vide supra).

Figure 3.12. Room temperature two-point I-V measurements of an unannealed WSe
2
film using a
0.85 mm 5.87 mm 75 nm channel.
75

The room temperature two-point I-V data of the unannealed films were collected using
two types of devices with different channel dimensions to verify the measured
conductivity values. Both unannealed and annealed films exhibited linear I-V behavior,
representative of ohmic contact between the WSe
2
nanosheets and the Al electrodes.
Eight devices from three different batches of WSe
2
nanosheets synthesized using
identical conditions gave an average conductivity of 0.6 ± 0.4 S cm
–1
, which is in the
same order of magnitude as WSe
2
single crystals (0.7 S cm
–1
)
20
and highly conductive
WSe
2
thin films (0.1 S cm
–1
).
24
Annealing the devices (475 °C) increased the average
conductivity by two orders of magnitude to 92 ± 27 S cm
–1
, similar to results published
for highly conductive MoS
2
films (100 S cm
–1
).
41
3.3. Conclusions and Future Work
In summary, a high-yielding synthesis of colloidal 2H-WSe
2
nanosheets was reported,
and shown to be phase pure by XRD. Preferential sheet alignment along the [001]
direction upon solution casting was reflected in the XRD pattern and the correspondingly
high conductivity values of the resulting thin films. Two-point conductivity
measurements for unannealed and annealed devices gave an average value of 0.6 and 92
S cm
–1
, respectively. Future work will focus on examining the viability of this material
for use in solution-processed, nanosheet-based devices.
76

3.4. Experimental Details
3.4.1. General Considerations
Tungsten(IV) chloride (WCl
4
, Strem, 97%), selenium powder (Alfa Aesar, 99.5%),
trichloroethylene (TCE, Alfa Aesar, 99.5+%), tetra-n-octylammonium bromide (TOAB,
Alfa Aesar, 98+%), 1,2-dichlorobenzene (DCB, Sigma-Aldrich, 99%), and tungsten(IV)
selenide (WSe
2
, Alfa Aesar, 99.8%) were all purchased and used without further
purification. Both dodecylamine (Alfa Aesar, 98+%), dried from CaO, and
tetramethylurea (TMU, Alfa Aesar, 99%) were distilled prior to use. Syntheses were
performed using standard Schlenk techniques under nitrogen in the absence of water and
oxygen.
3.4.2. Synthesis of WSe
2
nanosheets
In a typical synthesis, WCl
4
(150 mg, 0.46 mmol) was added to a three-neck round-
bottom flask fitted with a reflux condenser, stir bar, and rubber septum inside a glove
box. Deoxygenated, distilled dodecylamine (25 mL) was added to WCl
4
and then
dissolved under nitrogen by alternating between heating and an ultrasonic bath (5-10
min); care was taken as to not overheat the reaction mixture as WCl
4
is known to be
easily reduced in the presence of organics.
42
The dark bluish-black solution was then
heated to 95 °C and degassed for 3 minutes, cycling between vacuum and nitrogen three
times to eliminate adventitious water and dissolved oxygen. The reaction was further
heated to 150 °C (10 °C min
-1
), and
t
Bu
2
Se
2
(185 µL, 0.92 mmol) was quickly injected to
77

the reaction flask under nitrogen. The synthesis of
t
Bu
2
Se
2
was carried out according to
an improved version of a previously published method.
34,43
The temperature was
increased to 225 °C (10 °C min
–1
) and held at this temperature for 6 h while stirring. The
reaction was stopped by removing the heating mantle and allowing the reaction to cool to
room temperature, at which point 10 mL of a 20 mg mL
–1
TOAB-DCM solution was
used to disperse the reaction mixture. After vigorous mixing and sonication, the reaction
mixture was divided into two centrifuge tubes, and ethanol (25 mL) was added. A black
precipitate was obtained after centrifugation (6000 rpm for 1 min).
Dispersion/precipitation was repeated three more times using (1) 5 mL of 20 mg mL
–1

TOAB in DCM and 5 mL ethanol; (2) 5 mL of 40 mg mL
–1
TOAB in TCE, left
overnight, and 5 mL ethanol was added the next day; (3) 5 mL toluene and 5 mL ethanol
to remove excess TOAB. The product was washed a total of 4 times and finally re-
dispersed in ca. 5 mL toluene. Visually smooth films were drop-cast using ca. 40 mg
mL
–1
suspensions of WSe
2
in toluene.
3.4.3. Structural Characterization
Powder X-ray diffraction (XRD). Conventional XRD patterns were collected using a Cu
Kα radiation source (λ = 1.5406 Å) on a Rigaku Ultima IV diffractometer. Diffraction
patterns were recorded at room temperature in the range of 10-80˚. Scanning electron
microscope energy dispersive X-ray spectroscopy (SEM-EDX). SEM-EDX spectra were
collected on a JEOL JSM-6610 scanning electron microscope operating at 5 kV and
equipped with an EDAX Apollo silicon drift detector (SDD). Multiple regions of a
78

sample deposited on a Si substrate were analyzed. Elemental analysis (ICP-AES).
Elemental analysis of W and Se were performed by inductively coupled plasma atomic
emission spectroscopy at Galbraith Laboratories (Knoxville, TN). Raman spectroscopy.
Raman spectra were recorded under ambient conditions using a Horiba Jobin Yvon,
Xplora Raman Microscope System. An excitation source of 532 nm from a diode laser
was employed at a power level of 13.7 mW. Transmission electron microscopy (TEM)
and selected area electron diffraction (SAED). TEM and SAED were performed on a
JEOL JEM-2100 microscope at an operating voltage of 200 kV (or 100 kV for low
resolution TEM images), equipped with a Gatan Orius CCD camera. Samples were
prepared from dilute dispersions in toluene or TMU and deposited onto 300 mesh
Formvar-coated copper grids (Ted Pella, Inc.).
3.4.4. Device Characterization
Current-voltage measurements. The current dependence on applied test voltage
measurements were performed in air at room temperature using a Keithley 2420
SourceMeter (sensitivity = 100 pA). Data was collected at an interval of 40 mV. The
room temperature two-point I-V characteristics of the unannealed WSe
2
films were taken
using two types of devices with different channel dimensions (0.85 mm 5.87 mm 75
nm and 0.63 cm 2.51 cm 150 nm) to verify the conductivity values obtained. The
annealed WSe
2
films required a modified structure in which the material was annealed
(475 °C) first and then the Al electrodes were vapor deposited on top of the WSe
2
to
prevent contact failures. Assuming a uniform current flow along the film, we calculated
79

the resistance of the film to be the inverse of the slope of the I-V curve. The channel
conductivity (σ) for the unannealed films was estimated by assuming that the effective
height of the channel was given by the Al electrodes (σ = L/R A, where L is the channel
length, R is the measured resistance, and A the surface area of the Al electrodes facing
each other). Profilometry. The electrode thickness for the unannealed devices was
determined using a Sloan Dektak IIA profilometer. The channel conductivity (σ) for the
annealed films was estimated in a similar manner by calculating the effective height of
the channel by determining the concentration (20 mg mL
–1
) and depositing a known
volume of the WSe
2
sample in a 0.63 cm 2.51 cm channel. A density of 9.32 g cm
–3

for WSe
2
was used to derive channel thickness.
3.5.Acknowledgements
Acknowledgement is made to P. Erwin and S. Rodney (Thompson lab), J. Liu (Zhou lab),
H. Mahalingam (Steier lab), Dr. E. Couderc (Bradforth/Brutchey lab), Dr. C. Beier
(Brutchey lab), Dr. F. Rabuffetti (Brutchey lab), and Prof. Mark Thompson for both
experimental assistance and helpful discussions.
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photo-
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, and MoTe
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.” Phys.
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x
Ge
1-x
Se nanocrystals.” Chem. Mater. 2012, 24, 3514.
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84

Chapter 4. Low Temperature Solution-Phase Deposition of
SnSe Nanocrystalline Thin Films*
* Unpublished results

4.1. Introduction
The search for photovoltaic (PV) materials that can allow solar cells to compete with the
low cost of carbon-based fuels has propelled much work in the use of colloidal
nanocrystals as solar absorbers.
1-5
The IV-VI Sn and Ge monochalcogenide
semiconductors are an attractive alternative to more widely studied nanocrystalline PV
materials that contain toxic cadmium or lead, and to those that contain less abundant
elements like indium or gallium.
6,7
Tin selenide is a promising IV-VI PV material
because its direct (E
g dir
= 1.3 eV) and indirect (E
g ind
= 0.9 eV) bulk band gaps fall within
the ideal range for single junction solar cells.
8,9
Bulk SnSe is a p-type semiconductor that
exhibits a high hole mobility of 10
3
cm
2
V
–1
s
–1
,
10,11
and a high absorption coefficient of
5.5 × 10
5
cm
-1
.
12
In addition, SnSe is air stable and composed of earth abundant
elements. The use of colloidal SnSe nanocrystals in PVs is particularly appealing
because the material can be easily synthesized at much lower temperatures when
compared to the bulk, the synthesis can be easily scaled, and the band gap can be tuned
with particle size.
6
Furthermore, the organic ligands used to control nanocrystal size
allow for facile solution-based deposition and processing, an alternative to the energy-
intensive and vacuum-based chemical vapor deposition approaches.
2,13
Several groups
have reported on the colloidal synthesis and photovoltaic properties of SnSe. The first
85

example of well-defined and quantum-confined SnSe Nanocrystals was presented by
Franzman et al.
14
We obtained 19.0 ± 5.1 nm crystals that exhibited a blue-shifted direct
band gap of 1.71 eV. The SnSe nanocrystals were incorporated into a hybrid solar cell by
mixing them into a conductive polymer (MDMO-PPV) to act as the active layer. The
hybrid devices showed a short-circuit current density that was nearly twice that of the
neat polymer device, which translated into a power conversion efficiency enhancement of
100%. Baumgardner et al. further assessed the potential photovoltaic application of
colloidal SnSe Nanocrystals by reporting on the material’s photoconductive properties.
15

Quantum confinement effects were observed for the smaller SnSe nanocrystals (4 nm),
with an indirect band gap ranging from 1.2 eV to 0.9 eV, depending on nanocrystal size
(4 to 9 nm, respectively). The formic acid-treated SnSe nanocrystal films produced a
transient photocurrent response at a bias of 2 V µm
–1
under an illumination of 100 mW
cm
–2
. Vaughn II et al. also reported on the synthesis of photoconductive SnSe
nanosheets.
16
The resulting square-like nanosheets had uniform lateral dimensions of ca.
500 nm × 500 nm and a thickness that ranged between 10 and 40 nm. The indirect
optical band gap of E
g
= 1.0 eV did not reflect any quantum confinement effects, and no
photocurrent data was provided.
The deposition of colloidal PbSe nanocrystal thin films and the subsequent treatment with
simple amines or 1,2-ethanedithiol produces electronically coupled nanocrystal solids
that exhibit high carrier mobilities. Nozik et al. found that the amine treatments
preserved the size of the Nanocrystals, while removing much of the insulating native
oleate capping ligands in situ.
17
Despite the 2-7% removal of oleate via a
86

hydrazine/acetonitrile solution treatment, the resulting n-type transistors exhibited a high
mobility. In a similar manner, the group used a 1,2-ethanedithiol/acetonitrile solution to
remove the native ligands that originally dispersed the nanocrystals. The resulting field-
effect transistors were p-type and 30-60 times more conductive under illumination.
Ultimately, this procedure allowed for the fabrication of Schottky solar cells that
exhibited high current density values (J
SC
= 24.5 mA cm
–2
).
17c
Parkinson et al. used a
similar approach to treat their CZTS nanocrystals in situ, but instead of a solid state
device,
18
they evaluated the potential photovoltaic applicability of their thin films using a
photoelectrochemical method that has been used for rapid and nondestructive
assessments.
19-21
In addition, this method avoids the commonly observed electrical
shorting from a vapor-deposited metal back contact penetrating through the pores in the
film to the front contact in solid state devices. Following a similar approach, we report
on the photocurrent measurements of colloidal SnSe nanocrystal thin films. We compare
the material before and after the in situ treatment with hydrazine/acetonitrile or 1,2-
ethanedithiol/acetonitrile solutions.
4.2. Results and Discussion
4.2.1. Synthetic Methodology and Nanoparticle Characterization
The facile solution-phase synthesis of colloidal semiconductor nanocrystals has been
attained by our group using diorganodichalcogenides as effective chalcogen sources.
22-28

The synthesis of SnSe Nanocrystals was achieved by heating 0.56 g (3 mmol) tin(II)
chloride to 95 °C in dodecylamine (10 mL, 43 mmol) and dodecanethiol (2 mL, 8.4
87

mmol), followed by the injection of 280 µL
t
Bu
2
Se
2
(1.52 mmol). The reaction was
heated to 180 °C and quenched 25 min after nucleation using a water bath.

Figure 4.1. SnSe nanocrystals dispersed in toluene (left) settled after a few minutes, while those
dispersed in TMU (right) were more stable.

Figure 4.2. Thermal gravimetric analysis showing the decomposition endpoints of two different
SnSe dispersions. The improved workup procedure greatly reduced the mass loss to less than 4%
from ca. 14% (at 400 °C) using the original
14
workup procedure.

88

The product was thoroughly washed using DCM/EtOH, toluene/EtOH (twice),
tetramethylurea(TMU)/pentane, to then be re-disperse in toluene, for powder
characterizations; or TMU, for colloidal ink characterization (Figure 4.1). Thermal
gravimetric analysis (TGA) showed a mass loss of only 4% for the SnSe material using
the improved workup procedure, while the original workup procedure using DCM/ EtOH
and toluene/EtOH gave a mass loss of almost 14% (Figure 4.2).
The composition and structure of the SnSe Nanocrystals was confirmed via powder X-ray
diffraction (XRD) and energy-dispersive X-ray (SEM-EDX); while TEM images
revealed the expected rod-shaped particles (Figure 4.3). The powder X-ray (XRD)
diffraction pattern of the SnSe nanocrystals matched well that of PDF# 01-086-0232
(Figure 4.4).

Figure 4.3. TEM images of a typical batch of SnSe nanocrystals showing a rod-like shape.

89

Figure 4.4. Powder XRD pattern of SnSe Nanocrystals exhibiting an orthorhombic Pnma
structure (gray trace) and the matching PDF #01-086-0232 (black bars).

4.2.2. Deposition and Ligand Exchange Methodology
The SnSe nanocrystals were dispersed in tetramethylurea (TMU) at a concentration of ca.
50 mg mL
–1
for spin coating. An in situ ligand exchange was performed to remove the
insulating native ligands by soaking the thin films in a 1 M hydrazine (N
2
H
4
) or 0.1 M
1,2-ethanedithiol (edt) in dry acetonitrile using strict anaerobic conditions. Additional
layers were spin coated to increase the thickness of the thin films, and the ligand
exchange procedure carried out between each layer to minimize cracking and peeling.
FT-IR spectroscopy was used to compare the organic content before and after in situ
ligand exchange with edt and hydrazine. Figure 4.5 shows the strong ν(C–H) and ν(N–H)
stretching bands above 2750 cm
–1
. Although the ν(C–H) stretch decreases with
hydrazine treatment, the amine stretch above 3000 cm
-1
does not decrease, as there is no
90

amine displacement. Figure 4.6 shows how both the ν(C–H) and ν(N–H) stretching
bands between 3000 cm
–1
and 2750 cm
–1
decrease in intensity with edt. In the case of the
edt, the partial displacement of amine native ligands with edt could explain the decrease
in the amine stretch bands.

Figure 4.5. FT-IR staggered spectra of a SnSe nanocrystal thin film after consecutive
hydrazine/acetonitrile treatments.

Figure 4.6. FT-IR staggered spectra of a SnSe nanocrystal thin film after consecutive 1,2-
ethanedithiol/acetonitrile treatments.
91

4.2.3. Optical Thin Film Characterization
Scanning electron microscopy (SEM) was used to compare the film quality and coverage.
SEM imaging shows that the SnSe/TMU inks using the extended workup procedure
produced higher qu ality thin films than those obtained using the original workup
procedure (Figure 4.7).

Figure 4.7. SEM images of the SnSe/TMU derived films using the original workup procedure
(left) and an improved, extended work up procedure (right). Image on right shows a scratch on
the right side. Bottom SEM shows the film after hydrazine treatment with a dark scratch on the
left side.

92

The film conserved its integrity and was devoid of pin holes even after the hydrazine
ligand exchange with the hydrazine soak (Figure 4.7 bottom). Atomic force spectroscopy
(AFM) confirmed the SEM results, as the RMS roughness value (11 nm) remained
constant before and after hydrazine treatment.

Figure 4.8. Absorption spectrum of a SnSe/TMU derived film using the extended work up
procedure before and after hydrazine treatment.

UV-vis absorption spectroscopy was used to compare the spectra of films after prolonged
ligand exchange treatments to assess changes in band gap values. The broad absorption
spectrum of the as-made SnSe thin film exhibited very little change after a 15 h hydrazine
soak, which indicated that a shorter treatment would have a less pronounced effect on
film thickness (Figure 4.8). In addition, the slight increase in thin film absorption after
hydrazine treatment is consistent with literature results, which suggest that the decreased
interparticle spacing created a more densely packed film; this was corroborated using
93

variable angle spectroscopic ellipsometry, which showed a decrease in film thickness
after hydrazine treatment from 78.3 ± 0.3 nm to 63 ± 0.3 nm. The slight decrease in
absorption observed after the edt treatment suggests that some SnSe material was washed
away during the extended 15 h treatment (Figure 4.8). The direct band gap was
calculated (E
g
= 1.6 eV) derived using a Tauc Plot, and it did not significantly change
after the hydrazine treatment.

Figure 4.9. Absorption spectrum of a SnSe/TMU derived film using the extended work up
procedure before and after EDT treatment.

4.2.4. Photocurrent Response Measurements
The SnSe nanocrystal thin films’ potential for photovoltaic applications was assessed by
conducting photocurrent response measurements. Figures 4.10 and 4.11 show the
photocurrent results of a SnSe electrode made by spin coating two layers of the SnSe ink
onto ITO covered glass, which were then soaked in hydrazine or EDT (vide supra). The
current-voltage (I-V) curves show the p-type behavior exhibited by both non- and ligand-
94

exchanged films, as indicated by the increased photocurrent response as negative
potential is increased.
Prieto2
A photoelectrochemical cell containing 0.1 M
tetrabutylammonium hexafluorophosphate (TBAPf), a platinum counter electrode, a
silver wire reference electrode, and the SnSe thin films on ITO coated glass as the
working electrode. Two light emitting diodes were used for the incident 472 nm light
illumination, these were held on and off at 20 s intervals.

Figure 4.10. SnSe Nanocrystals spin coated on ITO with a constant applied potential, 472 nm
illumination, before and after hydrazine treatment.

The nanocrystalline SnSe films showed an average photocurrent density of 0.15 ± 0.03
µA cm
–2
which increased over 4 times to 0.64 ± 0.03 µA cm
–2
after the hydrazine in situ
ligand exchange. The introduction of ca. 1 mM Eu(NO
3
)
3
as a redox mediator in a 0.1 M
KCl aqueous solution, gave an average photocurrent value of 1.25 ± 0.01 µA cm
–2
before
and 1.85 ± 0.05 µA cm
–2
after the hydrazine ligand exchange. The use of ca. 1 mM Eu
triflate as a redox couple in a 0.1 M TBAP solution in dry acetonitrile gave an even larger
photocurrent response, but there was no difference between the as-made and the
95

hydrazine treated films. A control using a thin film made from native ligand SnSe should
give a more representative comparison of the effectiveness of both the hydrazine and edt
treatments.

Figure 4.11. SnSe Nanocrystals spin coated on ITO with a constant applied potential, 472 nm
illumination, before and after edt treatment.

4.3. Conclusion and Future Work
In summary, we report on the facile preparation of SnSe nanocrystal ink for the
preparation of thin film electrodes, and on the in situ ligand exchange of SnSe insulating
native ligands using hydrazine and 1,2-ethanedithiol. There was no significant change in
the absorption spectrum of the material or its band gap. We demonstrate control over
thin film morphology using SEM with a constant RMS roughness value of 11 nm.
Photocurrent response measurements showed a consistent p-type behavior across aqueous
and non-aqueous systems, while the use of an in situ ligand exchange holds promise in
the application of SnSe Nanocrystals for thin film based devices. Future work will
examine the incorporation of SnSe
96

4.4. Acknowledgements
Acknowledgement is made to Jose J. Araujo, Dr. Sean T. Roberts, Dr. David H. Webber,
Jannise J. Buckley, Matthew Greaney, Prof. Mark E. Thompson, Andrew Bartinsky,
Saptaparna Das, and Prof. Stephen E. Bradforth for both experimental assistance and
helpful discussions.
4.5. References
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2. Habas, S. E.; Platt, H. A. S.; van Hest, M. F. A. M.; Ginley, D. S. "Low-cost
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3. Guo, QJ ; Ford, GM; Agrawal, R; Hillhouse, HW. "Ink formulation and low-
temperature incorporation of sodium to yield 12% efficient Cu(In,Ga)(S,Se)
2
solar
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4. Hillhouse, H. W.; Beard, M. C. "Solar cells from colloidal nanocrystals:
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5. Guo, Q.; Hillhouse, H. W.; Agrawal, R. "Synthesis of Cu
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ZnSnS
4
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9. Shockley, W.; Queisser, H. J., "Detailed balance limit of efficiency of p-n junction
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10. Umeda, J. "Electrical properties of Sb-doped n-type SnSe." J. Phys. Soc. Jpn. 1961,
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11. Maier, H.; Daniel, D. R., "SnSe single-crystals - sublimation growth, deviation from
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12. Pejova, B.; Grozdanov, I. "Chemical synthesis, structural and optical properties of
quantum sized semiconducting tin(II) selenide in thin film form." Thin Solid Films
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13. Murray C. B.; Kagan C. R.; Bawendi M. G. “Synthesis and characterization of
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14. Franzman, M. A.; Schlenker, C. W.; Thompson, M. E.; Brutchey, R. L. "Solution-
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132, 4060.

15. Baumgardner, W. J.; Choi, J. J.; Lim, Y. F.; Hanrath, T. "SnSe nanocrystals:
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16. Vaughn II, D. D.; In, S. -I.; Schaak, R. E. "A precursor-limited nanoparticle
coalescence pathway for tuning the thickness of laterally-uniform colloidal
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17. a) Luther, J. M.; Law, M.; Song, Q.; Perkins, C. L.; Beard, M. C.; Nozik, A. J.
"Structural, optical, and electrical properties of self-assembled films of Pbe
nanocrystals treated with 1,2-ethanedithiol." ACS Nano, 2008, 2, 271. b) Law, M.;
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Beard, M. C.; Song, Q.; Reese, M. O.; Ellingson, R. J.; Nozik, A. J. "Schottky solar
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18. Riha, S. C.; Fredrick, S. J.; Sambur, J. B.; Liu, Y.; Prieto, A. L.; Parkinson, B. A.
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2
ZnSnS
4

photocathodes." ACS Appl. Mater. Interfaces 2011, 3, 58.

19. Scragg, J. J.; Dale, P. J.; Peter, L. M.; Zoppi, G.; Forbes, I. “New routes to
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2
ZnSnS
4
as an alternative absorber
material.” Phys. Stat. Sol. (B). 2008, 245, 1772.
20. Ye, H.; Park, H. S.; Akhavan, V. A.; Goodfellow, B. W.; Panthani, M. G.; Korgel,
B. A.; Bard, A. J. “Photoelectrochemical characterization of CuInSe
2
and Cu(In
1-
x
Ga
x
)Se
2
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using thermally labile tert-butylthiol for improved photocurrent in nanocrystal
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22. Franzman, M. A.; Perez, V.; Brutchey, R. L. “Peroxide-mediated synthesis of
indium oxide nanocrystals at low temperatures.” J. Phys. Chem. C 2009, 113, 630.
23. Webber, D. H.; Brutchey, R. L. “Photolytic preparation of tellurium nanorods.”
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2
nanocrystals.” Chem. Mater. 2010, 22, 1613.
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Ge
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99

Chapter 5. Solution-Phase Synthesis of NiS Nanocrystals*
*Unpublished results

5.1. Introduction
Millerite nickel sulfide (NiS) has been identified as binary semiconductor comprised of
earth abundant elements that could be used as alternative absorbing material in solar
cells;
1,2
and in general, various phases of nick el sulfide have also been explored as
electrode materials in lithium ion batteries,
3-5
counter electrodes in dye sensitized solar
cells,
6,7
and other catalytic applications such as H
2
evolution.
8
The synthesis of these
materials in the nanoscale regime would allow for the tailoring of particle size and band
gap,
9,10
but in the case of lithium ion batteries, the increased surface area of
nanostructured nickel sulfides could potentially decrease the diffusion length of lithium
ions and improve the performance of Li ion batteries.
3,11,12

Traditional methods to synthesize metal sulfides include reacting hydrogen sulfide gas
with an aqueous solution of the metal salt or by the high temperature heating of the
elemental sources in an evacuated silica tube.
13-15
In the case of NiS, Ni powder (reduced
with H
2
, 527 °C) and S powder were reacted in an evacuated silica tube for 7 d (727 °C);
while Ni
9
S
8
, first synthesized in 1969, was obtained by first reacting Ni sponge (reduced
with H
2
, 900 °C) and S in an evacuated sealed silica tube for 9 d (504 °C) to obtain α-
Ni
7
S
6
, which was then annealed in a sealed silica tube for 49 d (297 °C) to give the final
Ni
9
S
8
crystals.
16
In 2002, Qian et al. synthesized millerite NiS nanorods using a
100

hydrothermal approach reacting nickel chloride and sulfur or thiourea in ethylenediamine
and hydrazine hydrate in an autoclave (170 °C) for 12 h.
17
XRD results showed that a
mixture of NiS and NiS
2
was obtained when no hydrazine was used. TEM images
showed rod morphologies that are ~ 2 µm in length and 50 nm in diameter.

In 2011, Zhu
et al. synthesized both NiS nanorod- and nanoparticle-based hierarchical hollow
microspheres via a hydrothermal method in an alkaline solution of Na
2
S Ni(OH)
2
. In
turn, Ni(OH)
2
was synthezised by reacting Ni(NO
3
)
2
·6H
2
O, glycine and Na
2
SO
4
in an
alkaline solution at ca. 150 °C for 24 h.
18
XRD diffraction patterns showed a
combination of α-NiS and β-NiS, while the microspheres ranged between 1 and 5
micrometers in diameter by TEM; these were composed of either rods (~ 500 nm long, 50
nm wide) or ill-defined nanoparticles that seem to be around 100 nm in size. In 2014,
Aso and Tatsumisago synthesized both NiS and Ni
9
S
8
particles via thermal
decomposition of nickel acetylacetonate in a mixed solution of 1-dodecanethiol and
oleylamine or octadecene.
11,12
Oleyalmine was used as a coordinating solvent at different
temperatures, and regardless of reaction time, Ni
9
S
8
was the resulting phase of the
product. The use of octadecene as a noncoordinating solvent resulted in the formation of
NiS after heating Ni
9
S
8
at 280 °C for 5 h. Herein, we report on the facile synthesis of
well-defined NiS and Ni
9
S
8
nanoparticles in dodecylamine and dodecanethiol without the
need of a second annealing step.
101

5.2. Experimental Details
5.2.1. General Considerations
Nickel(II) chloride (NiCl
2
, Alfa Aesar, 98%), nickel(II) iodide (NiI
2
, Alfa Aesar, 99.5%,
anhydrous), N,N’-di-n-butylthiourea (Alfa Aesar, 98%), 1-dodecanethiol (Aldrich,
98+%), dibenzylamine (Alfa Aesar 98%), and tetra-n-octylammonium bromide (TOAB,
Alfa Aesar, 98+%) were used as received. Dodecylamine (Alfa Aesar, 98+%) was
distilled from CaO. Nanocrystal syntheses were performed under nitrogen, in the
absence of water and oxygen, using Schlenk techniques.
5.2.2. Nanocrystal Syntheses
A typical synthesis entailed the addition of NiCl
2
(0.1 g, 0.8 mmol) or NiI
2
(0.06 g, 0.2
mmol) into a three-neck round-bottom flask fitted with a reflux condenser, stir bar, and
rubber septum. Dodecylamine (5 mL) was injected into the reaction flask and then
stirred and heated to 95 °C to then cycle between vacuum and nitrogen three times. The
temperature was then increased to 175 °C (10 °C min
-1
), at which point a solution of
N,N’-di-n-butylthiourea (0.30 g, 1.6 mmol for NiCl
2
; 0.11 g , 0.6 mmol for NiI
2
) in
dibenzylamine (4 mL) was quickly injected into the system. In order to control particle
size and morphology, 1-dodecanethiol was injected ca. 30 s after the sulfur precursor.
The reaction mixture was further heated to 225 °C for 4 h in the case of the reaction with
NiCl
2
and 30 min for the reaction with NiI
2
. After cooling to room temperature, the
work-up procedure was promptly performed to avoid the solidification of dodecylamine
102

using a solution (50 mg/mL) of tetra-n-octylammonium bromide (TOAB) in DCM (5
mL) to dissolve the product; ethanol, along with light sonication and centrifugation (6000
rpm for 3 min), was used to precipitate the particles to yield either a black solid (β-NiS)
or a brassy colored solid (Ni
9
S
8
). The dispersion and precipitation procedure was
repeated twice with toluene (5 mL) and ethanol (5 mL) to yield the final product, which
was re-dispersed in toluene or other organic solvents for further characterization. Upon
drying, the NiS product exhibited a matte black color, while Ni
9
S
8
exhibited a metallic
brassy color.
5.2.3. Nanocrystal Characterization
The powder X-ray diffraction (XRD) analysis was performed on a Rigaku Ultima IV X-
ray diffractometer using a Cu Kα radiation source (λ = 1.54 A). 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.
5.3. Results and Discussion
5.3.1. NiS Micro- and Nanocrystals
Highly crystalline millerite NiS crystals were synthesized by the injection of 0.30 g (1.6
mmol) N,N’-di-n-butylthiourea dissolved in 4 mL dibenzylamine into a solution of 0.1 g
(0.8 mmol) NiCl
2
in 5 mL dodecylamine at 175 °C; the solution was then heated to 225
°C and held at this temperature for 4 h prior to quenching. Phase pure millerite crystals
were also obtained by injecting 0.4 g (2 mmol) N,N’-di-n-butylthiourea dissolved in 4 mL
103

dibenzylamine into a solution of 0.06 g (0.2 mmol) NiI
2
in 5 mL dodecylamine at 175 °C;
as before, the solution was heated to 225 °C, but it was instead allowed to react for 30
min. Tetra-n-octylammonium bromide (TOAB) was added during the work-up to
prevent agglomeration of the nanoparticles, following an adapted version of a previously
published method.
19
The final washed product was not highly dispersible in toluene, but
facile slow evaporation deposition resulted in visually smooth films. The omission of
TOAB during the work-up procedure yielded cracked films.
5.3.1.1. Structural Characterization

Figure 5.1. XRD pattern of β-NiS (millerite) using NiCl
2
, with N,N’-di-n-butylthiourea injected
at 175 °C and reacted for 4 h to 225 °C.

Powder X-ray diffraction (XRD) patterns of as-synthesized NiS match the rhombohedral
β-NiS (millerite) phase without any other crystalline nickel sulfide phases. Figure 5.1
104

and 5.2 display the XRD patterns of the NiS product matching the NiS phase (JCPDS no.
01-074-7239). Figure 5.1 corresponds to the NiCl
2
reaction, while Figure 5.2 is derived
from the reaction of NiI
2
. The careful addition of 1-dodecanethiol allowed for the
prevention of microscopic agglomerates.

Figure 5.2. XRD pattern of β-NiS (millerite) using NiI
2
, with N,N’-di-n-butylthiourea injected
at 175 °C and heated for 30 min to 215 °C.

Transmission electron microscope (TEM) analysis revealed the urchin-like morphology
of the NiS crystals with a slightly modified procedure using 1-dodecanethiol in the
reaction with NiI
2
(vide infra). Figure 5.3 shows the NiS nanorods arranged in urchin-
like systems that are well dispersed over large areas of the TEM grids; these samples
were diluted and drop-cast from toluene, while those dried from tetramethylurea showed
pronounced agglomeration under TEM. In an effort to reduce particle size and control
morphology, 1-dodecanethiol (2 mL) was injected after the sulfur source (N,N’-di-n-
105

butylthiourea, 10 eq) was added to NiI
2
(vide supra). The mixture was set to react for 45
min. The XRD diffraction peaks in fact widened, but prevented us from ascertaining the
absence of some other nickel sulfide phases.

Figure 5.3. TEM images of NiS dried from toluene showing the urchin-like structures.

5.3.2. Ni
9
S
8
Nanocrystals
Ni
9
S
8
nanocrystals were also synthesized using a solution of N,N’-di-n-butylthiourea
(0.11 g, 0.60 mmol) in dibenzylamine (2 mL) as the chalcogen source, which was
injected into a solution of NiI
2
(0.06 g, 0.2 mmol) in dodecylamine (5 mL). Excess 1-
106

dodecanethiol (4 ml) was injected at 30 s intervals in 0.5 mL increments. The mixture
was allowed to react to 225 °C for 45 min. Care was taken to avoid the solidification of
the dodecylamine/product mixture by dissolving in TOAB/DCM (vide supra), then 5 mL
toluene and precipitating with an equal volume of ethanol, light sonication, and
centrifugation (6000 rpm, 3 minutes). A dark solid was obtained by decanting the
supernatant. Dispersion and precipitation was repeated with toluene (5 mL) and ethanol
(5 mL) to yield the final product.
5.3.2.1. Structural Characterization
Figure 5.4. XRD diffraction pattern of Ni
9
S
8
crystals

Figure 5.4 shows the XRD pattern of the as-synthesized Ni
9
S
8
crystals with a close match
the godlevskite phase of Ni
9
S
8
; however, the presence of additional phases could not be
ruled out because of the diffraction peak breadth of the nanocrystals. TEM shows well-
dispersed crystals that are between 4-8 nm in diameter interspersed with larger (ca. 100
107

nm) particles (Figure 5.5 top), while the reaction using slightly different 1-dodecanethiol
(5 mL) amount and rates (1 mL/30 s) showed much less polydispersity (Figure 5.5
bottom).

Figure 5.5. TEM images of highly polydisperse Ni
9
S
8
crystals (top) and well-defined Ni
9
S
8

nanorods (bottom).

5.4. Conclusion
In summary, the use of N,N’-di-n-butylthiourea was used for the synthesis of NiS and
Ni
9
S
8
. The method is the foundation for the synthesis of potentially useful materials.
The method yielded phase pure β-NiS (millerite) and Ni
9
S
8
by the careful addition of 1-
108

dodecanethiol, which in turn helped control particle size and morphology. These
syntheses demonstrate how organochalcogenide precursors can be used for the synthesis
of various semiconductors.
5.5. Acknowledgements
Acknowledgement is made to Jose J. Araujo, Dr. David H. Webber, Ernesto Barron, and
Prof. Ilya Zharov for both experimental assistance and helpful discussions.
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Abstract (if available)

Abstract The solution phase deposition of inorganic semiconductors is crucial for the commercial‐scale fabrication of thin‐film based devices such as photovoltaics, thermoelectrics, and field effect transistors. The need for a truly dissolved inorganic semiconductor is highlighted by inexpensive high throughput processes such as spray coating and roll‐to‐roll printing that allow for large area and flexible substrate deposition, which are especially desirable for photovoltaic device commercialization. We have addressed the challenge of creating inks of several inorganic semiconductors via two low temperature techniques: 1) the dissolution and deposition of bulk tin sulfide (SnS) thin films using a hydrazine‐free solvent mixture and 2) the solution‐phase synthesis of tungsten selenide (WSe₂), tin selenide (SnSe), and nickel sulfide (NiS) nanoparticles. The indirect band gaps of SnS, WSe₂, and SnSe (1.1 eV, 1.2 eV, and 0.9 eV, respectively) make them attractive for light harvesting applications. In addition, SnS has been reported as a highly efficient photoelectrode material for water splitting, while WSe₂ has been incorporated into highly efficient photoelectrochemical devices (PCE = 17%). The methods we have developed allow for the facile solution‐phase deposition of films by spin‐coating for SnS and SnSe or by evaporation for WSe₂ and NiS. The materials are shown to be phase pure via powder XRD, SEM-EDS, diffuse reflectance UV-vis-NIR, Raman, and high resolution TEM

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Creator Antunez, Priscilla D. (author)

Core Title Solution‐phase synthesis and deposition of earth‐abundant metal chalcogenide semiconductors

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 08/21/2014

Defense Date 08/20/2014

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

Tag inorganic semiconductor,IV-VI semiconductor,nanocrystal,nanocrystals,nickel sulfide,NiS,OAI-PMH Harvest,photoresponse,SnS,SnSe,thin films,tin selenide,tin sulfide,tungsten diselenide,WSe₂

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Brutchey, Richard L. (committee chair), El-Naggar, Mohamed Y. (committee member), Thompson, Mark E. (committee member)

Creator Email antunez.priscilla@gmail.com,pantunez@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c3-463030

Unique identifier UC11286548

Identifier etd-AntunezPri-2836.pdf (filename),usctheses-c3-463030 (legacy record id)

Legacy Identifier etd-AntunezPri-2836.pdf

Dmrecord 463030

Document Type Dissertation

Format application/pdf (imt)

Rights Antunez, Priscilla D.

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

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