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Two-dimensional metal dithiolene metal-organic frameworks as conductive materials for solar energy conversion

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Two-dimensional metal dithiolene metal-organic frameworks as conductive materials for solar energy conversion

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Two-Dimensional Metal Dithiolene Metal-Organic Frameworks as
Conductive Materials for Solar Energy Conversion

by

Andrew James Clough

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 2019

ii

For Mom, Dad, and Rachel

iii

ACKNOWLEDGEMENTS

I must begin by thanking my advisor, Professor Smaranda C. Marinescu, without whose mentoring
I would be half the scientist I am today. It is amazing to think back on the six years that have
passed since those first electrochemistry experiments in Harry Gray’s wetbox on the third floor of
Noyes, and I am grateful for all I have learned over that time. I want to thank the other members
of my committee, Professors Brent Melot, Richard Brutchery, Sri Narayan, and Jaykanth
Ravichandran, for their scientific advice and professional insights throughout my time as a
graduate student. I am also grateful to my undergraduate advisor, Professor Karl O. Christe, and
master’s advisor, Professor Xianhui Bu, for their guidance and support through some of my most
transformative years as a young scientist.

To my lab mates in the Marinescu laboratory, your support has meant everything over the past six
years. To the other members of the original four Marinescu students, Courtney, Alon, and Damir,
having such talented and dedicated colleagues to embark with in a brand-new lab was an absolute
joy, and I couldn’t have asked for better people to have accompany me on this journey. To Eric,
thank you for always thinking about ways to improve yourself, me, and the Marinescu laboratory;
your relentless drive for research was an inspiration to me. To Nick, thank you for coordinating
the lab and being our primary method of scheduling Smaranda. To Ashley, Keying and Geo, thanks
for all the laughs, Love Letter, and general good times. To Jeremy, thank you for always looking
to improve yourself as a scientist; aiding in your development as a scientist has been a pleasure,
and I wish you the best of luck as the new research assistant for CNI’s XPS. To Colin, I wish you
the best of luck on the MOF project and look forward to your graduate career. To Stephanie, thanks
for your dedication to the BHT project (hopefully that paper will be published one day). To Ashley
de la Rosa, thanks for all of your hard work on the CoTHT jars, you did a great job. To Joseph,
you were the best undergraduate student a young graduate student could ask for. Thank you for
everything you contributed to the development of my project, and I wish you the best of luck in
medical school (just remember to come back and clean all that black crap off of the roof of my
fume hood!).

To my many collaborators, though there are too many to list here, thank you for your contributions
to my project. Many of our collaborations gave us key insights into the chemistry of my
iv

frameworks and provided several potential avenues to explore in the future. And of a course, a
special thank you to Professor Brent Melot, Abbey Neer and JoAnna Milam-Guerrero for teaching
me how to responsibly work with Ursula, the unruly PPMS.

To the staff of the Center for Electron Microscopy and Microanalysis (now the Core Center for
Excellence in Nano Imaging), especially John and Matt, thank you for providing a welcoming
environment for me to become an expert on the XPS. Your professional mentoring and support
when interacting with XPS users was much appreciated. To Matt Greaney, thank you for training
me well on the Axis Ultra, and showing me the ropes when I first took over from you.

To the Mount Whitney crew, John, Matt, Kyle, Amanda, Betsy, Kim: thank you for coming with
me on one of the most memorable (and grossest) treks I’ve ever been a part of. And to Betsy, an
extra special thanks for forcing me to drink water (DRINK!) and making sure I came back down
the mountain in one piece.

To the movie/game night crew, Abbi, Galactus, Chris, Chris, Leon, Val, Abby, and Danielle:
though I was frequently absent, being able to hang out with you all when I was having rough
patches really made those tough patches easier for me. I’m so lucky to be able to call you all my
friends.

To Keith, Rochelle, and Kyle/Jimmy/David: thank you for making our apartment the best, most
supportive living environment I could’ve asked for during my graduate studies. I will always
cherish our long nights of Secret Hitler, our varied and stimulating conversations, and for
apartment #29 being a welcome refuge from the trials of my graduate program.

And thank you to my sister, my father, and my mother: you have been the most loving,
supportive family for my entire life, and I could not have done this without you. Thank you for
always being there for me, for adding diversions when graduate life became too much (go
Knights go!), and for always believing in me, even when I didn’t believe in myself.
v

TABLE OF CONTENTS

Acknowledgements ......................................................................................................................... iii
Table of Contents .............................................................................................................................v
List of Figures ............................................................................................................................... vii
List of Schemes ............................................................................................................................. xiii
List of Tables ................................................................................................................................ xiv
Chapter 1. General Introduction ...................................................................................................1
1.1. General Background .....................................................................................................2
1.2. Conductivity in Metal-Organic Frameworks (MOFs) ..................................................4
1.3. Resistivity Measurements .............................................................................................6
1.4. Outline of this Work .....................................................................................................7
1.5. References .....................................................................................................................8
Chapter 2. Two-dimensional Metal-Organic Frameworks for Efficient Hydrogen
Evolution from Acidic Water ..........................................................................................15
2.1. Introduction .................................................................................................................16
2.2. Results and Discussion ...............................................................................................17
2.3. Conclusions .................................................................................................................43
2.4. Experimental Details ...................................................................................................44
2.4.1. General Considerations ................................................................................44
2.4.2. Synthesis of 1 ...............................................................................................45
2.4.3. Synthesis of 2 ...............................................................................................45
2.4.4. Electrochemical Methods.............................................................................48
2.5 References ....................................................................................................................49
Chapter 3. Metallic Conductivity in a Two-Dimensional Cobalt Dithiolene
Metal−Organic Framework ............................................................................................53
3.1. Introduction .................................................................................................................54
3.2. Results and Discussion ...............................................................................................55
3.3. Conclusions .................................................................................................................71
3.4. Experimental Details ...................................................................................................72
3.4.1. General Considerations ................................................................................72
3.4.2. Synthesis of CoTHT ....................................................................................72
3.4.3. Computational Modeling .............................................................................75
3.5. References ...................................................................................................................88
Chapter 4. Inducing Room Temperature Metallic Conductivity in a MOF via Oxidation ...92
4.1. Introduction .................................................................................................................93
4.2. Results and Discussion ...............................................................................................96
4.2.1. Synthesis and Characterization ....................................................................96
4.2.2. Density Functional Theory (DFT) Calculations ........................................101
4.2.3. X-ray Photoelectron Spectroscopy (XPS) and Magnetism Studies ...........103
4.2.4. Resistivity Studies ......................................................................................108
4.3. Conclusions ...............................................................................................................115
4.4. Experimental Details .................................................................................................116
4.4.1. General Considerations ..............................................................................116
4.4.2. Synthesis of FeTHT ...................................................................................116
vi

4.4.3. Computational Modeling ...........................................................................120
4.5. References .................................................................................................................134
Chapter 5. Conductivity in a Series of Metal–Benzenehexathiolate Two-Dimensional
Coordination Polymers ..................................................................................................140
5.1. Introduction ...............................................................................................................141
5.2. Results and Discussion .............................................................................................143
5.3. Conclusions ...............................................................................................................156
5.4. Experimental Details .................................................................................................156
5.4.1. General Considerations ..............................................................................156
5.4.2. Synthesis of FeBHT ..................................................................................157
5.4.3. Synthesis of CoBHT..................................................................................157
5.4.4. Synthesis of NiBHT ..................................................................................157
5.5. References .................................................................................................................158
Bibliography ................................................................................................................................163

vii

LIST OF FIGURES

Figure 2.1 Powder X-ray diffraction (PXRD) pattern of MOS 2 .................................................19
Figure 2.2 UV-Vis spectra of MOS 1 and 2 .................................................................................19
Figure 2.3 FTIR of [Co(bdt)2][nBu4N], benzenehexathiol, MOS 1, triphenylene-
2,3,6,7,10,11-hexathiol, and MOS 2 ..................................................................................20
Figure 2.4 High-resolution XPS spectra of MOS 1 on a glassy carbon plate ...............................21
Figure 2.5 High-resolution XPS spectra of MOS 2 grown using method B on a glassy
carbon plate ........................................................................................................................22
Figure 2.6 High-resolution XPS spectra of MOS 2 grown using method C on a glassy
carbon plate ........................................................................................................................23
Figure 2.7 SEM image, TEM SAED pattern, and SEM-EDS elemental maps of MOS 1
covering HOPG ..................................................................................................................24
Figure 2.8 TEM image and the select area electron diffraction pattern of MOS 1 on a
silicon nitride membrane....................................................................................................24
Figure 2.9 Polarization curves of MOS 1 and 2 in pH 10.0, 7.1, 4.4, 2.6, and 1.3
solutions .............................................................................................................................26
Figure 2.10 Scan rate dependence experiments for MOS 1 in pH 10.0 solution ..........................27
Figure 2.11 Scan rate dependence experiments for MOS 1 in pH 7.1 solution ............................28
Figure 2.12 Polarization curves of MOS 1 in aqueous solutions with pHs of 10.0, 7.1,
4.4, 2.6, and 1.0 ..................................................................................................................29
Figure 2.13 Polarization curves of MOS 2 in aqueous solutions with pHs of 10.0, 7.1,
4.4, and 2.6 .........................................................................................................................30
Figure 2.14 Polarization curves of MOS 1 and MOS 2 grown by method B in aqueous
solutions with pHs of 10.0, 7.1, 4.4, 2.6, and 1.0 ..............................................................31
Figure 2.15 Polarization curves of MOS 2 grown by method C in aqueous solutions with
pHs of 10.0, 7.1, 4.4, 2.6, and 1.3 ......................................................................................32
Figure 2.16 Current densities of MOS 1 as a function of the surface catalyst
concentration ......................................................................................................................33
Figure 2.17 Current densities of MOS 2 as a function of the surface catalyst
concentration. .....................................................................................................................34
Figure 2.18 Tafel plots of MOS 1 in aqueous solutions at pH 4.2 and pH 2.6 .............................34
Figure 2.19 Tafel plots of MOS 2 grown using method B in aqueous solutions at pH 4.4
and pH 2.6 ..........................................................................................................................35

Figure 2.20 Polarization curves of MOS 1 and MOS 2, BHTNa6, BHT, and
[Co(MeCN)6][BF4]2 at pH 1.3 ...........................................................................................35
Figure 2.21 Polarization curves of MOS 1 and dropcast [Co(bdt)2][nBu4N] in aqueous
solutions at pH 7.1, 4.4, 3.2, and 0.8..................................................................................36
viii

Figure 2.22 2 hour controlled potential electrolysis of MOS 1 and blank GCEs at
-0.55 V versus SHE in pH 2.6 and 1.0 ...............................................................................37
Figure 2.23 10 hour controlled potential electrolysis of MOS 1 and blank GCEs at
-0.55 V versus SHE in pH 2.6 and pH 1.0 .........................................................................38
Figure 2.24 10 hour controlled potential electrolysis of MOS 1 and blank GCEs at
-0.65 V versus SHE in pH 2.6............................................................................................38
Figure 2.25 XPS analysis of MOS 1 on a GCE after electrochemical studies. ............................39
Figure 2.26 XPS analysis of MOS 2 grown using method B on a GCE after
electrochemical studies ......................................................................................................40
Figure 2.27 XPS analysis of MOS 2 grown using the method C on a GCE after
electrochemical studies ......................................................................................................41
Figure 2.28 Polarization curves of [Co(bdt)2]
-
and a rinsed electrode in a 1:1
pH 1.3:acetonitrile solution ................................................................................................42
Figure 2.29 Image of the working compartment of the bulk electrolysis cell ..............................43
Figure 2.30 Picture of a typical set-up for the synthesis of the film (MOS 1) ..............................44
Figure 3.1 Spacefilling illustration of the fragment used to generate the model unit cell ............56
Figure 3.2 Model structures of the eclipsed pattern ......................................................................56
Figure 3.3 Experimental and simulated PXRD patterns of CoTHT ............................................57
Figure 3.4 Nitrogen sorption isotherms performed on CoTHT at 77 K.......................................58
Figure 3.5 Magnetic susceptibility, Curie-Weiss fit, and magnetization versus applied
magnetic field for CoTHT.................................................................................................58
Figure 3.6 Room temperature I-V trace of a CoTHT pressed pellet ............................................59
Figure 3.7 Variable-temperature resistivity data a pressed pellet and films of CoTHT ..............60
Figure 3.8 Arrhenius plot of the high temperature conductivity data of a pressed pellet of
CoTHT with Ea fit .............................................................................................................60
Figure 3.9 Overlay of the cooling and warming variable-temperature resistivity data for a
pressed pellet of CoTHT ...................................................................................................61
Figure 3.10 Synchrotron variable temperature PXRD patterns of solid CoTHT .........................61
Figure 3.11 Overlay of the variable temperature PXRD patterns of CoTHT focusing on
the [100] reflections ...........................................................................................................62
Figure 3.12 Overlay of the variable temperature PXRD patterns of CoTHT focusing on
the [001] reflections ...........................................................................................................62
Figure 3.13 SEM images of the cobalt dithiolene film CoTHT on glass a support .....................63
Figure 3.14 AFM studies of the cobalt dithiolene film CoTHT ..................................................63
Figure 3.15 AFM studies of the cobalt dithiolene film CoTHT ..................................................63
Figure 3.16 Images of the pellet or films of CoTHT and their puck assemblies for
conductivity measurements ................................................................................................64
Figure 3.17 Typical variable temperature I-V traces of the cobalt dithiolene film CoTHT ........64
Figure 3.18 Variable-temperature resistivity data for films CoTHT on glass supports ...............64
ix

Figure 3.19 Overlay of the warming and cooling variable-temperature resistivity data for
the cobalt dithiolene film of CoTHT.................................................................................65
Figure 3.20 Arrhenius plots of the high temperature conductivity data of the cobalt
dithiolene film CoTHT......................................................................................................65
Figure 3.21 Variable-temperature resistivity data for film CoTHT before and after a
two-hour exposure under vacuum at 90 °C........................................................................66
Figure 3.22 Arrhenius plot of the high temperature conductivity data of film CoTHT
before and after a two-hour exposure under vacuum at 90 °C ..........................................67
Figure 3.23 XPS data of the cobalt dithiolene film CoTHT collected at room
temperature before and after conductivity measurements .................................................68
Figure 3.24 Calculated electronic dispersion and density-of-states curve for CoTHT ................70
Figure 3.25 Contour map of the potential energy surface for offsets of alternate layers
along the a and b axes in CoTHT .....................................................................................71
Figure 3.26 Initial local magnetic moments and nominal total moments of seven
magnetic configurations tested for CoTHT ......................................................................77
Figure 3.27 Calculated spin density of the optimised model of CoTHT .....................................78
Figure 3.28 Calculated electronic dispersion and density-of-states curve for CoTHT ................80
Figure 3.29 Orbital-density plot showing the crystal orbitals associated with the partially-
occupied bands in the optimised model of CoTHT ..........................................................81
Figure 3.30 Calculated electronic dispersion and density-of-states curve for CoTHT with
a charge of –3 𝑒 per unit cell ..............................................................................................82
Figure 3.31 Energy change ∆𝐸 as a function of interlayer spacing (c-axis length) of
CoTHT ..............................................................................................................................83
Figure 3.32 Contour plot of the total energy of a bilayer of CoTHT as a function of layer
offset along the crystallographic a and b axes ...................................................................84
Figure 3.33 Calculated electronic density-of-states curves of CoTHT with different layer
spacings relative to the equilibrium value 𝑐 0
.....................................................................86
Figure 3.34 Calculated electronic density-of-states curves of a bilayer model of CoTHT
with several displacements of the layers along the a and b axis relative to each
other ...................................................................................................................................87
Figure 4.1 Synthesis, PXRD, and structure of the FeTHT framework ........................................97
Figure 4.2 Spacefilling model of the fragment used to generate the unit cell of FeTHT ............98
Figure 4.3 SEM images of FeTHT ...............................................................................................98
Figure 4.4 Nitrogen isotherms performed on the as-prepared FeTHT at 77 K ............................98
Figure 4.5 Variable temperature PXRD patterns of an oxidized FeTHT sample exposed
to air for 3 days at 60 °C ....................................................................................................99
Figure 4.6 Variable temperature PXRD patterns of an oxidized FeTHT sample exposed
to air for 3 days at 60 °C showing the [100] and [001] reflections ....................................99
Figure 4.7 PXRD pattern of an oxidized FeTHT sample exposed to air for 3 days at
60 °C. ...............................................................................................................................100
x

Figure 4.8 AFM studies of an oxidized FeTHT sample exposed to air for 3 days at 60 °C ......100
Figure 4.9 Nitrogen isotherms performed on an oxidized FeTHT sample exposed to air
for 3 days at 60 °C ...........................................................................................................101
Figure 4.10 Calculated band dispersion and electronic density of states curves for the
FeTHT framework with no guest species .......................................................................102
Figure 4.11 Calculated potential-energy surface associated with layer offsets (stacking
faults) in the pristine FeTHT framework ........................................................................103
Figure 4.12 Fitting of the XPS spectra of pristine FeTHT and samples exposed to
ambient atmosphere for 7 days ........................................................................................104
Figure 4.13 XPS spectra of the [Febdt2](HNEt3)2 molecular complex .......................................105
Figure 4.14 XPS spectra of the [Febdt2](HNEt3)2 molecular complex after 3 days of
oxidation in air .................................................................................................................105
Figure 4.15 XPS spectra of a [Pdbdt2]K2 molecular complex after 30 minutes of air
exposure ...........................................................................................................................106
Figure 4.16 XPS spectra from the oxidized [Pdbdt2]K2 species .................................................106
Figure 4.17 Magnetic hysteresis loops in pristine and 3 days air exposed FeTHT ...................108
Figure 4.18 Temperature-dependent susceptibility, Curie-Weiss fit, and hysteresis loop
of pristine FeTHT............................................................................................................108
Figure 4.19 Typical AFM studies of FeTHT films ....................................................................109
Figure 4.20 Variable-temperature I-V traces of FeTHT demonstrating Ohmic behavior .........109
Figure 4.21 Overlay of the temperature-dependent resistivity data for FeTHT films with
different thicknesses.........................................................................................................110
Figure 4.22 Arrhenius plots for FeTHT films ............................................................................110
Figure 4.23 Arrhenius plots using the variable range hopping model for FeTHT films ...........111
Figure 4.24 SEM images of FeTHT after resistivity studies .....................................................112
Figure 4.25 Variable-temperature PXRD patterns of pristine FeTHT .......................................113
Figure 4.26 Variable-temperature PXRD patterns of pristine FeTHT focusing on the
[100] and [001] reflections...............................................................................................113
Figure 4.27 Temperature-dependent resistivity data for an FeTHT sample with
increasing exposure to ambient atmosphere ....................................................................114
Figure 4.28 Temperature-dependent resistivity data for a FeTHT film as-prepared and
after 3 days of air exposure at 60°C .................................................................................115
Figure 4.29 Temperature-dependent resistivity data for a FeTHT film as-prepared and
after 3 days at 60°C under 1 atmosphere of helium. ........................................................115
Figure 4.30 Spin density of the lowest-energy magnetic configuration of the FeTHT
framework ........................................................................................................................123
Figure 4.31 Calculated band dispersion and electronic density of states curves for the
FeTHT framework ..........................................................................................................124
Figure 4.32 Orbital density showing states within 25 meV of the Fermi energy in the
FeTHT and the analogous CoTHT frameworks.............................................................125
xi

Figure 4.33 Calculated potential-energy surface associated with layer offsets in the
FeTHT framework ..........................................................................................................126
Figure 4.34 Illustration of the procedure for locating the minimum-energy c-axis length
(interlayer spacing) at a layer displacement of ∆a = 1.25 and ∆b = 1.25 Å ....................127
Figure 4.35 Optimized structures of bilayer models of FeTHT with layer displacements
corresponding to the local minima in Figure 4.32 ...........................................................128
Figure 4.36 Electronic density of states (DoS) curves of bilayer models of FeTHT in the
eclipsed configuration and staggered configurations .......................................................129
Figure 4.37 Band dispersion and electronic density of states (DoS) curves of the FeTHT
framework calculated with PBEsol and PBEsol+U with a Hubbard correction of
U = 5 eV applied to the Fe d states ..................................................................................132
Figure 4.38 Band dispersion and electronic density of states (DoS) curves of the Co
analogue of the FeTHT framework calculated with PBEsol and PBEsol+U with a
Hubbard correction of U = 5 eV applied to the Fe d states..............................................133
Figure 5.1 Scanning electron microscopy images of FeBHT ....................................................143
Figure 5.2 Scanning electron microscopy images of CoBHT ....................................................143
Figure 5.3 Scanning electron microscopy images of NiBHT .....................................................144
Figure 5.4 X-ray photoelectron spectra of FeBHT .....................................................................145
Figure 5.5 X-ray photoelectron spectra of CoBHT ....................................................................145
Figure 5.6 X-ray photoelectron spectra of NiBHT .....................................................................146
Figure 5.7 Temperature-dependent susceptibility, Curie-Weiss fit, and hysteresis loop for
FeBHT .............................................................................................................................147
Figure 5.8 Temperature-dependent susceptibility, Curie-Weiss fit, and hysteresis loop for
CoBHT ............................................................................................................................147
Figure 5.9 Temperature-dependent susceptibility, Curie-Weiss fit, and hysteresis loop for
NiBHT .............................................................................................................................148
Figure 5.10 Temperature-dependent resistivity and Arrhenius plot for a thick sample of
FeBHT .............................................................................................................................149
Figure 5.11 Temperature-dependent resistivity and Arrhenius plot for a thin sample of
FeBHT .............................................................................................................................149
Figure 5.12 Atomic Force Microscopy images of an FeBHT film ............................................149
Figure 5.13 Temperature-dependent I-V traces for a FeBHT film ............................................150
Figure 5.14 Temperature-dependent resistivity and Arrhenius plot for a thick sample of
CoBHT ............................................................................................................................151
Figure 5.15 Temperature-dependent resistivity and Arrhenius plot for a thin sample of
CoBHT ............................................................................................................................151
Figure 5.16 Atomic Force Microscopy images of a CoBHT film..............................................151
Figure 5.17 Temperature-dependent I-V traces for a CoBHT film ............................................152
Figure 5.18 Temperature-dependent resistivity and Arrhenius plot for a thin sample of
NiBHT .............................................................................................................................152
xii

Figure 5.19 Atomic Force Microscopy images of a NiBHT film ..............................................153
Figure 5.20 Temperature-dependent I-V traces for a NiBHT film .............................................153
Figure 5.21 Comparison of Tmetallic for FeBHT, CoBHT, NiBHT ............................................153
Figure 5.22 X-ray photoelectron spectroscopy of FeBHT after conductivity studies ................154
Figure 5.23 X-ray photoelectron spectroscopy of CoBHT after conductivity studies ...............154
Figure 5.24 X-ray photoelectron spectroscopy of NiBHT after conductivity studies ................155
xiii

LIST OF SCHEMES

Scheme 1.1 Schematic of the typical geometry for four-point van der Pauw measurements .........6
Scheme 2.1 Synthesis of the cobalt dithiolene films, MOS 1 and MOS 2 ....................................18
Scheme 3.1 Structure of the 2D cobalt dithiolene framework [Co3[THT]2]
3-
...............................55
Scheme 5.1 Structure of the 2D metal benzenehexathiolate films (M = Fe, Co, Ni) ..................142

xiv

LIST OF TABLES

Table 2.1. Current densities of MOS 1, MOS 2 and soluble [Co(bdt)2][nBu4N] at –0.63 V
versus SHE. ........................................................................................................................36
Table 3.1 Selected parameters for the variable-temperature resistivity analyses performed
on a pressed pellet or films of CoTHT. .............................................................................66
Table 4.1 Peak parameters to fit Fe
2+
and Fe
3+
multiplets in FeTHT.........................................104
Table 4.2 Summary of electrical transport data for FeTHT films ..............................................112
Table 4.3 Transport data for an FeTHT sample after exposure to ambient atmosphere for
varying times ....................................................................................................................113
Table 4.4 Optimized lattice parameters, total magnetic moments and total energies of
three trial magnetic supercells of the FeTHT framework ...............................................122
Table 4.5 Optimized lattice parameters, total magnetic moments and total energies of the
FeTHT framework with eclipsed layers and the three different staggered layer
configurations (stacking faults) ........................................................................................127
Tabe 4.6 Calculated lattice parameters and magnetic moments of the FeTHT framework
structure obtained using the PBEsol-D3 functional with and without a Hubbard
correction of U = 5 eV applied to the Fe d states.............................................................131
Table 4.7 Calculated lattice parameters and magnetic moments of the CoTHT analogue
of the FeTHT framework obtained using the PBEsol-D3 functional with and
without a Hubbard correction of U = 5 eV applied to the Co d states .............................133
Table 5.1 Electrical transport behavior for FeBHT, CoBHT, and NiBHT ...............................156

CHAPTER 1

General Introduction

2

1.1 General Background

Anthropomorphic climate change is a serious challenge that is threatening to cause a rise in global
temperatures, which are predicted to lead to hotter, dryer climates, a global rise in sea level as
polar sea ice is melted, and more intense storms.
1
The most significant contributions to man-made
climate change arise from the combustion of fossil fuels, which lead to emissions of greenhouse
gas on the order of ten thousand million tons of carbon-containing emissions per annum.
2
Of these
combustion products, carbon dioxide is the dominant product.
2
As sunlight is reflected off the
Earth’s crust, greenhouse gasses like CO2 absorb this energy as heat, and are the principle
contributors to the warming of the planet.
3
The carbon dioxide molecule is very
thermodynamically stable and can exist in the atmosphere for hundreds of years. The dominant
mechanism by which CO2 is removed from the atmosphere, the dissolution of CO2 into the
warming oceans, is also harmful to the world’s oceans, resulting in ocean acidification which can
harm oceanic ecosystems and lead to extinction of species vital to the food chain.
4
Carbon dioxide
concentrations in the atmosphere have recently surpassed 400 parts per million, and this
continuously increasing trend has occurred since the industrial revolution of the 1800s.
1
As a
result, new solutions and technologies are needed to address this growing problem while also
meeting the ever-increasing global demand for energy.
5

The production of renewable energy using sunlight is a potential solution to this problem.
6,7
Over
the past several decades, the development of technologies which can generate electricity from
renewable energy sources, such as wind, tidal, and especially photovoltaic solar panels, has
resulted in a dramatic reduction in the associated costs for power from environmentally friendly
and sustainable sources.
8
This trend is expected to continue, with renewable energy sources
projected to see the largest growth as a percentage of total energy production.
9
However, the
intermittent nature of these green energy production methods means that an effective method of
storing this clean energy must be developed which is cost-effective, scalable, and efficient.
5,10,11

One attractive method to address this challenge is the storage of energy in chemical bonds.
7,10,12

Hydrogen is an attractive candidate for this method of energy storage due to its carbon-free
composition, its high gravimetric energy density, and its relative abundance in water.
7,12
Hydrogen
is currently produced industrially through the steam reforming of methane, which also produces
carbon monoxide (eq. 1).
3

CH
4
+ H
2
O
∆
→ CO + 3H
2
(1)
The carbon monoxide produced by this reaction can then be further reacted via the water-gas shift
reaction to yield a further equivalent of hydrogen and produce CO2 (eq. 2).

CO + H
2
O
∆
→ CO
2
+ H
2
(2)
This process thus produces hydrogen gas, which can be used as an energy carrier or for other
industrial processes; however, the co-production of CO2 means that an alternative, carbon-free
source of hydrogen is required. Water, with its two atoms of hydrogen per molecule, is a green
and abundant potential source of hydrogen.
7,13
Additionally, fuel cell technology already exists
which can cleanly convert H2 into water while generating electricity in the process.
14

Unfortunately, the conversion of H2O into H2 is energy-intensive, and there is an ongoing need for
materials which can efficiently, reliably, and cheaply perform these conversions on a large scale.
14

The overall equation for the splitting of water into hydrogen and oxygen gas is given in equation
3.

2H
2
O

→ 2H
2
+ O
2
(3)
The hydrogen evolution reaction (HER) is the half-reaction which reduces the protons in water
using two electrons to form dihydrogen gas (equation 1.1).
2H
+
+ 2e
−

→ H
2
(4)
Platinum is the most efficient electrocatalyst for the hydrogen evolution reaction. However, due to
its scarcity and related high cost, catalysts based on scalable, earth-abundant metals are needed if
water splitting is to be a feasible strategy for solar energy conversion. Nature can accomplish the
splitting of water with incredible efficiency near the thermodynamic potential using base metals.
The active site of the [FeFe]-hydrogenase has evolved over billions of years to contain several key
motifs which enable its high catalysis using abundant metals; these include an open metal site, an
[4Fe4S] cluster which can store electrons, and pendant amines which facilitate proton transport.
15-
17
These design elements have influenced the rational study of man-made catalysts for hydrogen
evolution.

Homogeneous catalysts, those were the active species is in solution, are valuable for studying
catalytic mechanisms and rationally improving their performance. A number of complexes have
4

been studied for their HER properties.
18-20
Cobaloximes have been shown to reduce protons to
form dihydrogen gas under photoelectrochemical conditions.
20,21
A nickel catalyst with phosphine-
amine arms (a P2N2 catalyst) mimics the pendant proton relays of the [FeFe]-hydrogenase active
site; this enables it to perform efficient catalysis at over 100,000 turnovers per second in wet
acetonitrile.
22,23
And a cobalt dithiolene catalyst operates at low overpotentials in acetonitrile:water
solutions, demonstrating stability and high activity in aqueous media.
24-26

Heterogeneous catalysts are insoluble or deposited materials which function on an electrode
surface; these catalysts can be extremely active at low overpotentials and effective as catalysts for
solar energy conversion.
27
A number of different compositions of nanoparticles have been
investigated for hydrogen evolution, including nickel phosphide nanoparticles.
28-31
However, their
density of active sites can be lower than their bulk loading would suggest. An example of this can
be found with MoS2, which has been shown to only be active along its edge sites.
30,32

Understanding the mechanism by which these heterogeneous catalysts operate is also more
challenging than their homogeneous counterparts.

Efforts have been made to combine the advantages of both homogeneous and heterogeneous
catalysis by immobilizing homogeneous catalysts on electrode surface. These surface attachment
techniques include covalent attachment
33,34
and non-covalent adsorption
35-37
. Examples of these
include the tethering of Nickel P2N2 catalysts to carbon nanotube surfaces
33
, the adsorption of
cobalt dithiolene complexes to graphene and other graphitic supports,
35,38
the use of pyrene groups
to adsorb nickel P2N2 catalysts
37
, and the co-adsorption of rhenium aminothiolato-based
complexes to graphitic supports
39
. While these methods enable the immobilization of
electrocatalysts on conductive supports, their catalyst loading concentrations are limited to
monolayers, which limits their performance. Additionally, the adsorbed molecular species can
desorb from the surface under catalytic conditions, which limits their stability and application in
water splitting devices.

1.2 Conductivity in Metal-Organic Frameworks (MOFs)

Metal-Organic frameworks (or MOFs), initially developed as synthetic zeolites with new
topologies, are typically synthesized by the solvothermal reaction of metal ions or nodes with
5

organic ligands, or linkers, to generate three-dimensional crystalline structures with a relatively
low defect concentration.
40-42
MOFs typically exhibit very high internal surface areas and
permanent porosity, properties with potential applications in gas storage, gas separation, and
catalysis.
43-47
Most MOFs are insulating in nature; this is due to the hard interactions between the
metal centers and the oxygen atoms in the carboxylate groups which are commonly used in linkers.
As a result, their band structures exhibit flat bands with very low band dispersion, and large gaps
between their valence and conduction bands.
48-50
To enable device applications where charge
transport in MOFs is necessary, several strategies have been employed to attempt to improve the
intrinsically poor conductivity of MOFs. These include using linkers with special properties or
traits, such as redox-active moieties, the use of dopants containing redox-active components
51
or
conductive species like nanoparticles
52
, incorporation of electrically conductive polymers,
53
and
the formation of composite materials with conductive materials like graphene
54
or carbon
composites.
55
However, these strategies can also limit the device applications of MOFs, and can
cause issues with the stability of the host MOF or occupy much of their internal surface area, which
limits their applications. Intrinsically conductive MOFs are thus a desirable target to overcome
these challenges.

Recently, MOFs have emerged in the literature which exhibit appreciable conductivity. A Ni/Cu-
containing 3D MOF using pyrazine and dithiolene linkages exhibited modest conductivity and
showed significant enhancements in its transport properties after treatment with I2 vapor.
56,57

Through-space conductivity pathways have also shown the ability to enhance charge transport in
porous 3D frameworks.
58,59
The incorporation of planar two-dimensional redox-active linkers has
led to dramatic breakthroughs in the conductivity of MOFs.
50,60-65
These systems frequently use
dithiolene
56,57,66-78
, diimine
60,79-82
, or semiquinones/cathecolate-based linkers
83-86
, and represent a
new class of lower-dimensionality MOFs with superior conductivity properties.
72,87
The record
holder for conductivity in a MOF is the copper benzenehexathiolate 2D coordination polymer,
which exhibits conductivity exceeding 2,500 S/cm, metallic character at all temperatures, and a
superconducting transition below 1 K.

6

1.3 Resistivity Measurements

In this work, we will be making use of resistance measurements to describe the intrinsic
conductivity of two-dimensional materials. At their most basic level, resistance measurements
typically measure the voltage with the presence of an applied current, or the current in the presence
of an applied voltage; this relationship is governed by Ohm’s law (equation 5). Although this work
will not use Ohm’s law to determine the resistance of samples directly, for reliable voltage
measurements, the system must display a linear current/voltage trend across the potential range
where data will be taken. Non-Ohmic contacts can occur because of a variety of causes, including
Schottky barriers, a p-n junction, a rectifying junction, and others causes.

V = IR (5)
The sheet resistance of a material can be determined using a number of methods. For the purposes
of this work, all the sheet resistances are determined using the van der Pauw method, which is a
four-contact method that enables sheet resistance determination on an arbitrary sample geometry.
The equations for this determination are given in equations 6 and 7.

Scheme 1.1. Schematic of the typical geometry for four-point van der Pauw measurements.

𝑅 𝐴 =
(
|𝑉 12,34
|
𝐼 +
|𝑉 34,12
|
𝐼 +
|𝑉 21,43
|
𝐼 +
|𝑉 43,21
|
𝐼 )
4
(6)
𝑅 𝐵 =
(
|𝑉 23,41
|
𝐼 +
|𝑉 41,23
|
𝐼 +
|𝑉 32,14
|
𝐼 +
|𝑉 14,32
|
𝐼 )
4
(7)
When RA and RB are equivalent, the sheet resistance can then be calculated by:
R
s
=
πR
ln2
(8)
7

The resistivity of the material, ρ, can thus be described by Rs divided by its thickness, t. The units
of this value are typically reported in Ohm centimeters, or Ω·cm.
ρ = R
s
∙ t (9)

1.4 Outline of this work

In this dissertation, we will look at a new method of immobilizing catalytically-active units by
incorporating them into a two-dimensional metal organic framework. Our approach used a
liquid:liquid interfacial reaction scheme to generate thin films of the material at the interface.
These films can then be isolated as powders or readily deposited on the desired substrate.

In chapter two, we will look at two novel two-dimensional metal-organic frameworks developed
using cobalt metal centers and benzenehexathiol or 2,3,6,7,10,11-triphenylenehexathiol ligands as
linkers. We examine the structure, morphology, and electrochemical behavior, observing a
reversible feature that indicates that redox chemistry is possible with these surface-immobilized
species; these studies confirm that the cobalt-dithiolene moieties are immobilized on the surface
and that they have sheet-like morphologies. We then examine them in acidic solutions and observe
that they efficiently catalyze the production of hydrogen gas from acidic water with good stability
and efficiencies. Their operating overpotentials are found to be 340 and 530 mV versus the
standard hydrogen electrode with 100% Faradaic efficiencies toward hydrogen production.

In chapter three, we will examine the conductivity of the CoTHT framework. In contrast with
most reported MOFs in the literature, CoTHT exhibits appreciable conductivity, with
semiconducting behavior at temperatures near room temperature. Despite the numerous
computational studies which suggest that many 2D MOFs should exhibit metallic character, there
were no experimental examples in the literature that demonstrated metallic conductivity in a MOF.
We found that at low temperatures the CoTHT MOF exhibits a complex transition to metallic
conductivity, a finding which was unprecedented at the time; there is now only one other MOF
known to possess metallic conductivity, a CuBHT-based 2D MOF which also exhibits
superconductivity at temperatures below 2 K.

8

In chapter four, we examine the conductivity of the structurally-analogous FeTHT two-
dimensional MOF. Because of a number of reports in the literature that suggested inducing mixed-
valency in iron-based MOFs could be a strategy for dramatically enhancing conductivity, we
investigated the transport properties of after it was subjected to aerobic oxidation. While the
conductivity decreased with increasing oxidation, the observed metallic transition temperature
increased, reaching near room temperature with sufficient oxidation.

In chapter five, we report initial studies on the conductivity of the MBHT series of 2D MOFs (M
= Fe, Co, Ni). We find that conductivity increases across the group (Fe < Co < Ni). Interestingly,
the FeBHT and CoBHT exhibit transitions to metallic behavior at low temperatures (and the
CoBHT only at cryogenic temperatures), whereas the NiBHT exhibits no such transition. The
morphology and surface chemistry are also discussed.

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CHAPTER 2

Two-Dimensional Metal-Organic Frameworks for Efficient Hydrogen Evolution
from Acidic Water

A portion of this chapter has appeared in print:

Clough, A. J.; Yoo, J.; Mecklenburg, M. H.; Marinescu, S. C. “Two-Dimensional Metal–Organic
Surfaces for Efficient Hydrogen Evolution from Water.” J. Am. Chem. Soc., 2015, 137, 118–121

16

2.1 Introduction

Energy harvested directly from sunlight offers a desirable approach toward fulfilling the global
need for clean energy.
1,2
Hydrogen produced through the reduction of water is an attractive
candidate as a clean and renewable fuel. Much work has gone into developing homogeneous and
heterogeneous hydrogen evolving catalysts made from nonprecious metals for applications in
efficient, scalable energy storage. A variety of base metals homogeneous
3-9
and
heterogeneous
10
hydrogen evolving catalysts have been developed recently. The construction of
efficient and practical devices for electrocatalytic splitting of water requires the attachment to
electrodes of hydrogen evolving catalysts based on nonprecious metals that can operate and are
robust under acidic aqueous conditions, under which proton exchange membrane-based
electrolysis is operational.
11,12
A variety of cobalt species have been adsorbed onto the electrodes
by controlled potential electrolysis and display H2-evolving activity.
13-16
Molecular catalysts are
attractive because the ligand environment allows for tuning of their reduction potentials and
chemical properties. However, the reported methods for the immobilization of molecular catalysts
onto electrodes are scarce. Demonstrated grafting methods include the covalent attachment of
nickel
11,17
or cobalt complexes
18
onto carbon-based supports (multiwalled carbon nanotubes or
glassy carbon), the coordination of cobaloximes to vinylpyridine GaP surfaces,
19,20
and the
immobilization of molecular catalysts via pyrene groups.
21,22
However, these methods suffer from
low coverage; the maximum catalyst loading achieved is 10
–8
molcatalyst/cm
2
.
22
The attachment of
catalysts to surfaces in a well-defined fashion remains a great challenge. Incorporation of active
sites in an extended catalytic surface was envisioned to maintain the properties of the molecular
catalysts such as activity and reduction potential, while also increasing the catalyst loading and
rendering the material more robust and transferable to an electrode surface by direct deposition.

Two-dimensional (2D) covalent– or metal–organic frameworks (COFs/MOFs) afford control of
atomic layers deposited on surfaces.
23-29
These materials can be grown on a variety of supports
and are characterized by high charge carrier mobility and high surface-to-volume ratio. Moreover,
the heterogeneous nature of the COFs/MOFs allows for easy separation, reusability, and enhanced
stability, making them attractive in the context of developing integrated photoanode and
photocathode materials in a solar fuel cell. However, despite their great promise, the catalytic
properties of these nanoscopic architectures have not been described. We report here the
17

development of metal–organic surfaces (MOS) and their applications as efficient electrode
materials for the reduction of acidic water.

2.2 Results and Discussion

Cobalt dithiolene species are among the most efficient molecular catalysts for the hydrogen
evolution reaction (HER).
30-34
Related nickel dithiolenes species have been investigated as well
for HER;
35,36
however, their chemistry is under debate.
30,37
In order to access extended
architectures with integrated cobalt dithiolene catalytic sites, a trinucleating conjugated ligand,
benzenehexathiol (BHT), was employed in reactions with cobalt(II) through a liquid–liquid
interfacial process (Scheme 2.1). An acetonitrile/ethyl acetate solution of [Co(MeCN)6][BF4]2 was
gently layered on top of an aqueous solution of sodium benzenehexathiolate (C6S6Na6). The
organic solvents were allowed to evaporate over 1 h, leaving behind a black film (1) at the gas–
liquid interface. Film 1 can be deposited on a variety of supports, such as glassy carbon or highly
oriented pyrolytic graphite (HOPG), by immersing the support face down into the reaction mixture.
The generated MOS 1 was washed and dried. Additionally, the black powder was collected on a
fine porosity frit and washed with water and methanol. The measured elemental composition of
the black powder corresponds to a molecular formula of (CoC4S4Na)n. This method has been
extended to frameworks based on triphenylene-2,3,6,7,10,11-hexathiolate (2).
18

Scheme 2.1. Synthesis of the cobalt dithiolene films, MOS 1 and MOS 2. Reprinted with
permission from Ref. 118. Copyright 2015 American Chemical Society.

Powder X-ray diffraction (PXRD) studies of 2, obtained by method A, revealed a crystalline
structure with peaks at 2θ = 4.5, 9.1, 12.0 and 15.6°, indicative of long-range order within
the ab planes (Figure 2.1). The peak at 2θ = 4.5° corresponds to a distance of ∼20 Å, which is
equal to the value expected for the pore diameter of MOF 2 determined using known bond lengths
and angles for molecular cobalt dithiolene species. An additional weaker and broader peak at 2θ =
26.7° is indicative of poorer long-range order along the c direction, as expected for layered
materials. A similar PXRD pattern has been reported for layered MOF Ni3(2,3,6,7,10,11-
hexaiminotriphenylene)2.
29
UV–vis spectra of 1 and 2 display broad bands between 270 and 600
19

nm wavelengths (Figures 2.2), analogous with the UV–vis spectra of MOF Ni3(2,3,6,7,10,11-
hexaiminotriphenylene)2.
29
The FTIR spectra of 1 and 2 show the disappearance of the strong
signal present in the hexathiols at 2500 cm
–1
, which corresponds to the S–H stretching vibration
(Figure 2.3).

Figure 2.1. Powder X-ray diffraction (PXRD) pattern of 2. Reprinted with permission from Ref.
118. Copyright 2015 American Chemical Society.

Figure 2.2. UV-Vis spectra of MOS 1 (a) and 2 (b). Reprinted with permission from Ref. 118.
Copyright 2015 American Chemical Society.

20

Figure 2.3: FTIR spectra of [Co(bdt)2][nBu4N] (a, purple), benzenehexathiol (b, blue), MOS 1 (c,
green), terphenylene-2,3,6,7,10,11-hexathiol (d, orange), and MOS 2 (e, red). Reprinted with
permission from Ref. 118. Copyright 2015 American Chemical Society.

X-ray photoelectron spectroscopy (XPS) analysis of MOS 1 revealed the presence of Co, S, and
Na (Figure 2.4). Two sets of peaks are observed in the cobalt region, with binding energies of
∼780 and ∼795 eV, which correspond to the 2p3/2 and 2p1/2 levels in the expected 2:1 ratio.
Deconvolution of these signals generates four peaks; the peaks at 779.2 and 794.2 eV are assigned
to Co
III
, whereas the ones at 780.9 and 795.4 eV are assigned to Co
II
.
38
This assignment is in
agreement with the electronic structure of cobalt bisdithiolene complexes, which was interrogated
by a variety of spectroscopic and theoretical studies, and is best represented by the resonances
[Co
III
(bdt)2]
−1
↔ [Co
II
(bdt)(bdt
•
)]
−1
(bdt =1,2-benzenedithiolate).
39
Three additional peaks are
observed with binding energies of 1071.4, ∼228, and ∼163 eV, which correspond to Na 1s, S 2s,
and S 2p, respectively. The broad peaks at 230.0 and 166.0 eV are assigned to the shape-up
satellites, which are often observed in bisdithiolene complexes.
26,40,41
XPS analysis of
MOS 2 revealed similar peaks with the ones observed for 1 (Figures 2.5 and 2.6). These data
support assignment of the MOS with cobalt dithiolene moieties linked by hexathiolate nodes.
21

Figure 2.4. X-ray photoelectron spectroscopy (XPS) analysis of MOS 1 on glassy carbon plate.
(a) Co 2p core level XPS spectrum; (b) Na 1s core level XPS spectrum; (c) S 2s core level XPS
spectrum; (d) S 2p core level XPS spectrum. Reprinted with permission from Ref. 118. Copyright
2015 American Chemical Society.

22

Figure 2.5. X-ray photoelectron spectroscopy (XPS) analysis of MOS 2 grown by method B on a
glassy carbon plate electrode. (a) Co 2p core level XPS spectrum; (b) S 2s core level XPS
spectrum; (c) S 2p core level XPS spectrum. Reprinted with permission from Ref. 118. Copyright
2015 American Chemical Society.

23

Figure 2.6. X-ray photoelectron spectroscopy (XPS) analysis of MOS 2 grown by method C on a
glassy carbon plate electrode. (a) Co 2p core level XPS spectrum; (b) Na 1s core level XPS
spectrum; (c) S 2s core level XPS spectrum; (d) S 2p core level XPS spectrum. Reprinted with
permission from Ref. 118. Copyright 2015 American Chemical Society.

The coverage and thickness of the film on HOPG was evaluated by scanning electron-microscopy
(SEM). A top-down micrograph is illustrated in Figure 2.7a. Cross-sectional micrographs were
obtained after depositing a protective layer of Pt and milling the sample with a Ga
+
focused ion
beam (FIB). The cross section of the film revealed an average thickness of 360(40) nm. The
average surface catalyst concentration of MOS 1, determined using the electrochemical methods
described below, was 3.7(4) × 10
–6
molCo/cm
2
. Detection of the electron diffraction from MOS 1,
using select area electron diffraction (SAED) techniques, was inhibited by the film’s beam
sensitivity (the images lasted less than 200 ms). The electron diffraction pattern of MOS 1 on
silicon nitride membranes displays a hexagonal symmetry, suggesting that the material is weakly
crystalline (Figures 2.7b and 2.8). Similar patterns were detected at several locations in the sample.
Elemental mapping of cobalt and sulfur using SEM–energy dispersive spectroscopy (SEM–EDS)
are illustrated in Figure 2.7c-e, indicating that these elements are homogeneously distributed and
colocalized throughout the film.
24

Figure 2.7 (a) A typical SEM image of MOS 1 covering HOPG. (b) An SAED pattern of MOS 1
on silicon nitride membrane taken from ~ 200 nm diameter region of the material. The beam
sensitivity (peaks lasted less than 200 ms) prevented higher order diffraction peaks from being
acquired. (c) A secondary electron image of MOS 1; the brighter regions indicate the film, the
darker regions the HOPG support. Elemental mapping of cobalt (d) and sulfur (e) using SEM-
EDS. Reprinted with permission from Ref. 118. Copyright 2015 American Chemical Society.

Figure 2.8. (a) Characteristic TEM image of MOS 1 on silicon nitride membrane. (b) Select area
electron diffraction pattern from MOS 1. Radiolysis prevented the diffraction pattern from being
detected on our camera for longer than 200 ms. The lattice parameter, determined from nine
diffraction patterns, was 0.49 ± 0.03 nm. Reprinted with permission from Ref. 118. Copyright
2015 American Chemical Society.
25

The electrochemistry of MOS 1 was investigated by cyclic voltammetry (Figure 2.9a). When the
MOS-modified glassy carbon electrode was immersed in pH 10.0 buffer, a broad reversible redox
couple was observed, with an E° of −0.30 V vs SHE (−0.50 V vs SCE). The peak current was
directly proportional to the scan rate over the range of 1–100 mV/s (Figures 2.10 and 2.11). The
peak separation (ΔEp) is small (<20 mV) at scan rates below 10 mV/s, indicative of rapid electron
transfer between the glassy carbon electrode and MOS 1. At higher scan rates (10–100 mV/s), the
separation between the reduction and oxidation peaks increases with the scan rate. The observed
voltammetric profile is indicative of a surface-confined redox couple. The value of the redox
couple for MOS 1 is similar to the one reported for the molecular analogue, Co(bdt)2
–
, which was
observed at −0.64 V vs SCE in 1:1 H2O/MeCN solutions, and assigned to Co(bdt)2
–1/–
2
.
30,31,42
Maximum average surface catalyst concentrations are similar for MOS 1 and 2 (3.7(4) ×
10
–6
molCo/cm
2
for MOS 1, and 2.5(3) × 10
–6
molCo/cm
2
for MOS 2), as estimated from the
integration of the electrochemical wave at pH 10.0. Similar values are obtained from ICP–MS
measurements of the digested catalyst films, suggesting that the majority of the cobalt centers are
electrochemically active.
26

Figure 2.9. Electrochemical studies of MOS 1 and 2. (a) Polarization curves of MOS 1 (1.9(2) ×
10
–6
molCo/cm
2
) in 0.1 M NaClO4 aqueous solutions at pH 10.0 (red), 7.1 (blue), 4.4 (green), 2.6
(purple), and of blank glassy carbon electrode at pH 2.6 (dashed black); scan rate, 20 mV/s. (b)
Polarization curves of MOS 1 (red, 0.7(1) × 10
–6
molCo/cm
2
) and MOS 2 (blue, method B, 1.1(1) ×
10
–6
molCo/cm
2
) in pH 1.3 H2SO4solution; scan rate, 100 mV/s. Reprinted with permission from
Ref. 118. Copyright 2015 American Chemical Society.

27

Figure 2.10. Scan rate dependence experiments for MOS 1 (catalyst loading: 1.2(1)

´10
–6

molCo/cm
2
) in 0.1 M NaClO4 aqueous solutions at pH 10.0. (a) scan rates from 1 mV/s to 20 mV/s;
(b) scan rates from 25 mV/s to 100 mV/s; (c) peak reduction and oxidation current density for
MOS 1 as a function of the scan rate; (d) peak reduction and oxidation current density for MOS 1
as a function of the square root of the scan rate; (e) oxidation and reduction peak separation of
MOS 1 as a function of the scan rate. Reprinted with permission from Ref. 118. Copyright 2015
American Chemical Society.

28

Figure 2.11. Scan rate dependence experiments for MOS 1 (catalyst loading: 1.9(2)

´10
–6

molCo/cm
2
) in 0.1 M NaClO4 aqueous solutions at pH 7.1. (a) scan rates from 1 mV/s to 20 mV/s;
(b) scan rates from 20 mV/s to 900 mV/s; (c) Peak reduction and oxidation current density for
MOS 1 as a function of the scan rate; (d) Peak reduction and oxidation current density for MOS 1
as a function of the square root of the scan rate. Reprinted with permission from Ref. 118.
Copyright 2015 American Chemical Society.

As the pH of the aqueous solution is lowered, an increase in current is observed (Figures 2.9a and
2.12) indicating that a catalytic reaction is taking place. A peak-shaped catalytic current was
observed at −0.65 V vs SHE at pH 2.6 for MOS 1. A potential value of −0.51 V vs SHE (−0.35 V
vs RHE) was required to reach a current density of 5 mA/cm
2
at pH 2.6. By comparison,
unmodified glassy carbon electrodes display insignificant increase in current. Similar behavior
was observed for MOS 1 and MOS 2 prepared using method B (Figures 2.13 and 2.14) and MOS
2 prepared by method C (Figure 2.15). The solutions that resulted after the electrochemical studies
of MOS 1 at each pH were subjected to cyclic voltammetry in the presence of a blank glassy carbon
electrode. Negligible current densities were observed, indicating that the modified electrode does
not generate soluble materials responsible for catalysis. Additionally, ICP–MS measurements
indicate that the amount of cobalt present in solution is negligible.
29

Figure 2.12. Polarization curves of MOS 1 (catalyst loading: 1.9(2) ⨯ 10
-6
molCo/cm
2
) in 0.1 M
NaClO4 aqueous solutions at pH 10.0 (a), 7.1 (b), 4.4 (c), 2.6 (d), 1.0 (e); scan rates: 20 mV/s for
(a)–(d) and 100 mV/s for (e). Reprinted with permission from Ref. 118. Copyright 2015
American Chemical Society.

30

Figure 2.13. Polarization curves of MOS 2 (catalyst loading: 1.1(1) ⨯ 10
-6
molCo/cm
2
) grown by
method B in 0.1 M NaClO4 aqueous solutions at pH 10.0 (a), 7.1 (b), 4.4 (c), 2.6 (d), and 1.3 (e);
scan rates: 20 mV/s for (a)–(d) and 100 mV/s for (e). Reprinted with permission from Ref. 118.
Copyright 2015 American Chemical Society.

31

Figure 2.14. Polarization curves of MOS 1 (catalyst loading: 2.4(2) ⨯ 10
-6
molCo/cm
2
) and MOS
2 (catalyst loading: 0.6(1) ⨯ 10
-6
molCo/cm
2
) grown by method B on glassy carbon electrodes in
0.1 M NaClO4 aqueous solutions at pH 10.0 (a), 7.1 (b), 4.4 (c), 2.6 (d); scan rate: 20 mV/s.
Reprinted with permission from Ref. 118. Copyright 2015 American Chemical Society.

32

Figure 2.15. Polarization curves of MOS 2 (catalyst loading: 0.6(1) ⨯ 10
-6
molCo/cm
2
) grown by
method C in 0.1 M NaClO4 aqueous solutions at pH 10.0 (a), 7.1 (b), 4.4 (c), 2.6 (d), and 1.3 (e);
scan rates: 20 mV/s for (a)–(d) and 100 mV/s for (e). Reprinted with permission from Ref. 118.
Copyright 2015 American Chemical Society.

At the lowest pH value (pH = 1.3), and in the presence of MOS 1 or MOS 2, large current densities
can be observed (41 and 31 mA/cm
2
for MOS 1 and 2, respectively, at −0.8 V vs SHE) (Figure
2.9b). It is noteworthy that the voltammograms of the MOS 1 and 2 at pH 1.3 are distinct. The
measured onset for H2-evolution for MOS 1 is at −0.28 V vs SHE, whereas for MOS 2 is at −0.48
V vs SHE, suggesting that the active material is not identical for MOS 1 and 2. Catalytic current
33

densities of MOS 1 and 2, measured at potentials of −0.55 and −0.85 V vs SHE, respectively,
increase linearly with catalyst loading (Figures 2.16 and 2.17). Tafel analyses of MOS 1 and 2gave
Tafel slopes between 149 and 189 mV/dec, and exchange current densities of 10
–5.3(1)
A/cm
2
at pH
4.2 (Figures 2.18 and 2.19), which are comparable to the ones reported for cobalt complexes
grafted onto amine-modified multiwalled carbon nanotubes.
18
Glassy carbon electrodes modified
with starting materials only (cobalt(II) or thiolates) display insignificant increase in current (Figure
2.20). The activities of MOS 1 and 2 were compared to that of the molecular analogue, [Co(bdt)2]
(Figure 2.21 and Table 2.1. Glassy carbon electrodes modified with [Co(bdt)2]
−
by the drop-
casting method or in solution display a very small increase in current, suggesting that
immobilization via MOS provides an increase in activity and stability.

Figure 2.16. Current densities of MOS 1 measured at –0.55 V vs SHE and pHs of 4.2 (red) and
2.6 (blue) as a function of the surface catalyst concentration. The catalyst loading was determined
by integrating the peaks area at pH 10.0. Scan rate: 20 mV/s. Reprinted with permission from Ref.
118. Copyright 2015 American Chemical Society.

34

Figure 2.17. Current densities of MOS 2 measured at –0.85 V vs SHE and pH 1.3 as a function of
the surface catalyst concentration. The catalyst loading was determined by integrating the peaks
area at pH 10.0. Scan rate: 100 mV/s. Reprinted with permission from Ref. 118. Copyright 2015
American Chemical Society.

Figure 2.18. Tafel plots of MOS 1 (catalyst loading: 3.0(3)

´10
–7
molCo/cm
2
) in 0.1 M NaClO4
aqueous solutions at pH 4.2 (red; Tafel slope of 149 mV/dec) and pH 2.6 (blue; Tafel slope of
108 mV/dec); scan rate: 5 mV/s. Reprinted with permission from Ref. 118. Copyright 2015
American Chemical Society.

35

Figure 2.19. Tafel plots of MOS 2 (catalyst loading: 1.2(1)

´10
–6
molCo/cm
2
) grown by method
B in 0.1 M NaClO4 aqueous solutions at pH 4.4 (red; Tafel slope of 189 mV/dec; exchange
current density of 10
-5.2
A/cm
2
) and pH 2.6 (blue; Tafel slope of 161 mV/dec; exchange current
density of 10
-5.4
A/cm
2
); scan rate: 0.5 mV/s. Reprinted with permission from Ref. 118. Copyright
2015 American Chemical Society.

Figure 2.20. (a) Polarization curves of MOS 1 (catalyst loading: 1.6(2)

´10
–6
molCo/cm
2
),
dropcast sodium benzenehexathiolate (BHTNa6 – color), benzenehexathiol (BHT), and
[Co(MeCN)6][BF4]2 (color) in H2SO4 aqueous solutions at pH 1.3 and scan rates of 100 mV/s;
(b) Polarization curves of MOS 2 (catalyst loading: 1.6(2)

´10
–6
molCo/cm
2
), dropcast sodium
benzenehexathiolate (BHTNa6 – color), benzenehexathiol (BHT), and [Co(MeCN)6][BF4]2
(color) in H2SO4 aqueous solutions at pH 1.3 and scan rates of 100 mV/s. Reprinted with
permission from Ref. 118. Copyright 2015 American Chemical Society.

36

Figure 2.21. Polarization curves of MOS 1 (catalyst loading: 1.6(2)×10
–6
molCo/cm
2
) and
dropcast [Co(bdt)2][nBu4N] in 0.1 M NaClO4 aqueous solutions at pH 7.1 (a), 4.4 (b), 3.2 (c),
and 0.8 (d); scan rates: 20 mV/s for (a)–(c) and 500 mV/s for (d). Reprinted with permission from
Ref. 118. Copyright 2015 American Chemical Society.

Entry Compound Current Density (A/molCo) Medium
1 MOS 1 4.3(7) × 10
4
pH1.3
2 MOS 2 1.2(3) × 10
4
pH13
3 [Co(bdt)2]
-
1.4(2) × 10
2
1:1 pH 1.3:MeCN
Table 2.1. Current densities of MOS 1 (0.7(1) × 10
-6
molCo/cm
2
), MOS 2 (1.1(1) × 10
-6
molCo/cm
2
)
and soluble [Co(bdt)2][nBu4N] (0.3 mM) at –0.63 V vs. SHE. Reprinted with permission from Ref.
118. Copyright 2015 American Chemical Society.

Controlled potential electrolysis of MOS 1 on a glassy carbon electrode in 0.1 M NaClO4
citrate/phosphate buffer at pH 2.6 and −0.8 V vs SHE consumed 45 coulombs of charge after 2 h
(Figure 2.22). Analysis of the gas mixture in the headspace of the working compartment of the
electrolysis cell by gas chromatography confirmed production of H2 with a Faradaic yield of 97 ±
37

3%. In the presence of MOS 1 and a mixture of NaClO4/HClO4 at pH 1.0 and −0.8 V vs SHE, 57
coulombs of charge were consumed in 1 h. The durability of MOS 1 in pH 2.6 aqueous solution
was further assessed in a longer-duration controlled potential electrolysis experiment. MOS 1
affords a continuous increase in charge build-up over a 10 h controlled potential electrolysis at
−0.55 or −0.65 V vs SHE (Figures 2.23 and 2.24). As H2 bubbles generated during the 10 h
controlled potential electrolysis cause the peeling-off of the material, the decrease in catalyst
loading leads to some lowering in current flow. XPS analyses of MOS 1 and 2 after
electrochemical studies display peaks similar to the ones observed before electrolysis, suggesting
that the material is stable under reductive and acidic conditions (Figures 2.25–2.27). By
comparison, unmodified glassy carbon electrodes display insignificant H2-evolution activity.

Figure 2.22. Controlled potential electrolysis of MOS 1 (4.5(5) ⨯10
-7
molCo/cm
2
) at –0.8 V vs
SHE in 0.1 M NaClO4 aqueous solutions at pH 2.6 (blue), and 1.0 (red), and of blank glassy
carbon plate electrodes at pH 2.6 (green), and 1.0 (black dashed). Reprinted with permission
from Ref. 118. Copyright 2015 American Chemical Society.

38

Figure 2.23. Controlled potential electrolysis of MOS 1 (1.1(1) ⨯ 10
-6
molCo/cm
2
) at –0.55 V vs
SHE in 0.1 M NaClO4 aqueous solutions at pH 2.6 (red), and of blank glassy carbon plate
electrodes (black dashed). It is important to note that, for longer controlled potential electrolysis
measurements, it is difficult to maintain equal pressure on both chambers of the H-cell. This
leads to variances in the surface area of the working electrode immersed in solution, and
therefore variances in the measured current. The spikes in current are due to the formation of H2
bubbles, or due to a short pause in the experiment to re-equilibrate the levels of the liquids in the
two compartments of the H-cell. The catalyst loading of MOS 1 measured at the end of the
experiment is 0.4(1) ⨯ 10
-6
molCo/cm
2
. Reprinted with permission from Ref. 118. Copyright 2015
American Chemical Society.

Figure 2.24. Controlled potential electrolysis of MOS 1 (0.9(1) ⨯ 10
-6
molCo/cm
2
) at –0.65 V vs
SHE in 0.1 M NaClO4 aqueous solutions at pH 2.6 (red), and of blank glassy carbon plate
electrodes (black dashed). It is important to note that, for longer controlled potential electrolysis
measurements, it is difficult to maintain equal pressure on both chambers of the H-cell. This
leads to variances in the surface area of the working electrode immersed in solution, and
therefore variances in the measured current. The spikes in current are due to the formation of H2
bubbles, or due to a short pause in the experiment to re-equilibrate the levels of the liquids in the
two compartments of the H-cell. The catalyst loading of MOS 1 measured at the end of the
experiment is 0.4(1) ⨯ 10
-6
molCo/cm
2
. Reprinted with permission from Ref. 118. Copyright 2015
American Chemical Society.

39

Figure 2.25. X-ray photoelectron spectroscopy (XPS) analysis of MOS 1 on glassy carbon plate
electrode after electrochemical studies. (a) Co 2p core level XPS spectrum; (b) Na 1s core level
XPS spectrum; (c) S 2s core level XPS spectrum; (d) S 2p core level XPS spectrum. Reprinted
with permission from Ref. 118. Copyright 2015 American Chemical Society.

40

Figure 2.26. X-ray photoelectron spectroscopy (XPS) analysis of MOS 2 grown by method B on
glassy carbon plate electrode after electrochemical studies. (a) Co 2p core level XPS spectrum;
(b) S 2s core level XPS spectrum; (c) S 2p core level XPS spectrum. Reprinted with permission
from Ref. 118. Copyright 2015 American Chemical Society.

41

Figure 2.27. X-ray photoelectron spectroscopy (XPS) analysis of MOS 2 grown by method C on
glassy carbon plate electrode after electrochemical studies. (a) Co 2p core level XPS spectrum;
(b) Na 1s core level XPS spectrum; (c) S 2s core level XPS spectrum; (d) S 2p core level XPS
spectrum. Reprinted with permission from Ref. 118. Copyright 2015 American Chemical
Society.

To test whether the molecular cobalt dithiolene species decomposes during catalysis to generate a
heterogeneous material active for H2-evolution, the following control experiment was performed.
Cyclic voltammograms of a 0.3 mM solution of [Co(bdt)2]
−
in a 1:1 mixture of pH 1.3 aqueous
H2SO4 solution and 0.1 M KNO3 acetonitrile solution display a maximum current density of 6.0
mA/cm
2
at −0.63 V vs SHE (Figure 2.28). The electrode was then rinsed with water and MeCN
several times to remove any physisorbed complex and transferred to a fresh 1:1 mixture. The
observed current densities of the rinsed electrode were negligible, suggesting that [Co(bdt)2]
−
does
not decompose during catalysis to generate a heterogeneous material active for H2-evolution.
42

Figure 2.28. Polarization curves of 0.3 mM [Co(bdt)2]
-
in a 1:1 mixture of pH 1.3 aqueous
H2SO4 solution and 0.1 M KNO3 acetonitrile solution (blue) and of the rinsed electrode (5 ⨯ 5
mL water, followed by 5 ⨯ 5 mL acetonitrile; 10 min. each washing step) in a fresh 1:1 mixture
(red); scan rates: 100 mV/s. Reprinted with permission from Ref. 118. Copyright 2015 American
Chemical Society.

The materials prepared here operate under fully aqueous conditions. Overpotentials of 0.34 and
0.53 V are required for MOS 1 and 2, respectively, to reach current densities of 10 mA/cm
2
at pH
1.3. In comparison, other immobilized H2-evolving catalysts require much higher overpotentials
to reach current densities of 10 mA/cm
2
.
18
Similar overpotentials to the ones observed here have
been reported for molecular cobalt dithiolene complexes in mixtures of aqueous and organic
solvents suggesting that, although in an extended material, the environment around cobalt
maintains the properties of the molecular catalyst.
30,31
This feature is important for further rational
tuning of reactivity. The maximum average surface catalyst concentration observed for MOS 1 is
3.7(4) × 10
–6
molCo/cm
2
, which is 2 orders of magnitude higher than the maximum catalyst loading
reported.
22
This high surface concentration is indicative of a multilayered material. In addition to
the low overpotential, the high catalyst loading and activity, MOS 1 displays remarkable stability
under acidic conditions, and moderate durability in longer-duration controlled potential
electrolysis. Theoretical studies performed on the molecular cobalt dithiolene species suggest that
the mechanism of hydrogen evolution involves protonation of the sulfur sites on the dithiolene
ligands after the initial Co
III/II
reduction.
32,42
This may eventually lead to ligand loss and
decomposition in molecular systems. The high stability observed here for MOS 1 is ascribed to
43

the network environment. Enhanced photochemical hydrogen production was also reported for a
molecular diiron benzenedithiolate catalyst incorporated into a metal–organic framework.
43

2.3 Conclusions

Cobalt dithiolene species are among the most efficient molecular catalysts for hydrogen
evolution.
30-34
We demonstrate here the successful integration of the cobalt dithiolene catalysts
into MOS to give very active electrocatalytic cathode materials for hydrogen generation from fully
aqueous solutions. The materials generated display high catalyst loadings and remarkable stability
under acidic conditions. These results indicate that immobilization as MOS provides a significant
increase in activity and stability for these cobalt catalysts and thus paves the way toward
development of practical devices.

Figure 2.29. Image of the working compartment of the bulk electrolysis cell, which is separated
from the counter electrode by a fine porosity frit. The chamber contains a glassy carbon plate
electrode in 0.1 M NaClO4 aqueous solution at pH 2.4, and the reference electrode, placed in a
separate compartment connected by a Vycor tip. Reprinted with permission from Ref. 118.
Copyright 2015 American Chemical Society.

44

Figure 2.30. Picture of a typical set-up for the synthesis of the film (MOS 1). Reprinted with
permission from Ref. 118. Copyright 2015 American Chemical Society.

2.4 Experimental Details

2.4.1 General Considerations

All manipulations of air and moisture sensitive materials were conducted under a nitrogen
atmosphere in a Vacuum Atmospheres drybox or on a dual manifold Schlenk line. The glassware
was oven-dried prior to use. Acetonitrile and dichloromethane were degassed with nitrogen and
passed through activated alumina columns and stored over 4 Å Linde-type molecular sieves. Ethyl
acetate and ethanol were placed under vacuum and refilled with argon (10

´). Deuterated solvents
were dried over 4 Å Linde-type molecular sieves prior to use. Proton NMR spectra were acquired
at room temperature using Varian spectrometers and referenced to the residual
1
H resonances of
the deuterated solvent (
1
H: CDCl3, δ 7.26; C6D6, δ 7.16; CD2Cl2, δ 5.32) and are reported as parts
per million relative to tetramethylsilane. Elemental analyses were performed by Complete
Analysis Laboratories, Inc., Parsippany, New Jersey, 07054 or Robertson Microlit Laboratories,
1705 U.S. Highway 46, Suite 1D, Ledgewood, New Jersey, 07852. All the chemical regents were
purchased from commercial vendors and used without further purification. Benzenehexathiol
44

(BHT), terphenylene-2,3,6,7,10,11-hexathiol (THT)
45
, and [Co(bdt)2][nBu4N]
30
were prepared
45

according to the reported procedures. Water was deionized with the Millipore Milli-Q UF Plus
system (18.2 M·cm resistivity). All other chemical regents were purchased from commercial
vendors and used without further purification. The pHs of the aqueous solutions were measured
with a benchtop VWR Symphony pH meter.

2.4.2. Synthesis of 1

An aqueous solution of NaOH (0.5 ml, 0.753 M, 0.377 mmol) was added to a suspension of C6S6H6
(17.1 mg, 0.063 mmol) in water (25 mL). The reaction mixture was allowed to stir at room
temperature until the solids dissolved. The pale yellow solution was filtered if necessary and
transferred to a crystallizing dish (70 mm diameter

´ 50 mm height). An acetonitrile : ethyl acetate
1:4 (v:v) solution of [Co(MeCN)6][BF4]2 (7.5 mM, 1 mL) were gently added via a glass pipette to
the aqueous solution to cover ~80% of the surface area. The organic solvents were allowed to
evaporate over one hour at room temperature, leaving behind MOS 1 as a black film at the air-
liquid interface. The black film was deposited on the carbon-based supports – glassy carbon (GC)
or highly oriented pyrolytic graphite (HOPG) – by immersing the support face down into the
reaction mixture. The dimensions of the supports are: GC electrode – diameter = 3 mm; GC plate
– 5 cm

´ 1 cm

´ 0.3 cm; HOPG electrode – diameter = 9.3 mm; HOPG plate – 5 cm

´ 1 cm

´
0.3 cm. The volatiles were allowed to evaporate at room temperature or under vacuum. The
supported material was washed with water and methanol. Alternatively, the black film was
collected on a fine porosity frit and washed with water, ethanol, and dichloromethane. The black
powder was dried under vacuum. Anal. Calcd for [Co3(BHT)2][Na3] (Co3S12C12Na3): C, 18.60;
Co, 22.82; Na, 8.90; S, 49.67. Found: C, 18.54; Co, 22.69; Na, 8.78; S, 49.57.

2.4.3 Synthesis of 2

Method A.

A 120 ml jar was charged with a solution of CoCl2·6H2O (35.8 mg, 0.15 mmol) in water (20 mL).
An ethyl acetate solution (20 mL) was gently added via a glass pipette to the aqueous solution. N-
methyl-2-pyrrolidone (NMP) (0.2 mL, 2.07 mmol) was added to solid terphenylene-2,3,6,7,10,11-
hexathiol (THT) (17 mg, 0.041 mmol). The mixture was suspended in ethyl acetate (5 mL), sealed,
and sonicated for 45 minutes to disperse the powder. The cloudy light green suspension was then
46

diluted with ethyl acetate (15 mL) and added gently via a glass pipette to the jar. After 5 days, a
black film was observed at the liquid-liquid interface. The black film was washed with water and
methanol, and solvent exchanged for acetone (3

´ 20 mL). The recovered solids were heated to
100 °C under dynamic vacuum for 12 h. Anal. Calcd for [Co3(THT)2][H3]•MeCOMe•1/2H2O
(Co3C45H36S12O0.5): C, 44.66; Co, 14.61; S, 31.79. Found: C, 44.24; Co, 14.77; S, 31.82.

Method B.

N-methyl-2-pyrrolidone (NMP) (0.2 mL, 2.07 mmol) was added to a suspension of terphenylene-
2,3,6,7,10,11-hexathiol (THT) (4 mg, 0.01 mmol) in ethylacetate (5 mL). The slightly soluble
mixture was sealed in a 20 mL vial and sonicated for 45 minutes to disperse the powder. The
cloudy light green suspension was gently added via a glass pipette to a crystallizing dish (70 mm
diameter

´ 50 mm height) containing a solution of CoCl2·6H2O (21.5 mg, 0.09 mmol) in water
(20 mL). The organic solvents were allowed to evaporate over several hours at room temperature,
leaving behind 2 as a black film at the air-liquid interface. The black film was deposited on the
carbon-based supports – glassy carbon (GC) or highly oriented pyrolytic graphite (HOPG) – by
immersing the support face down into the reaction mixture. The dimensions of the supports are:
GC electrode – diameter = 3 mm; GC plate – 5 cm

´ 1 cm

´ 0.3 cm; HOPG electrode – diameter
= 9.3 mm; HOPG plate – 5 cm

´ 1 cm

´ 0.3 cm. The volatiles were allowed to evaporate at room
temperature or under vacuum. The supported material was washed with water and methanol.

XPS data were collected using a Surface Science Instruments M-Probe ESCA controlled by Hawk
Data Collection software (Service Physics, Bend OR; V7.04.04) or a Kratos AXIS Ultra
instrument. The monochromatic X-ray source was the Al K α line at 1486.6 eV, directed at 35
o
to
the sample surface (55
o
off normal). Emitted photoelectrons were collected at an angle of 35
o
with
respect to the sample surface (55
o
off normal) by a hemispherical analyzer. The angle between the
electron collection lens and X-ray source is 71
o
. Low-resolution survey spectra were acquired
between binding energies of 1–1100 eV. Higher-resolution detailed scans, with a resolution of
~0.8 eV, were collected on individual XPS lines of interest. The sample chamber was maintained
at < 2 ´ 10
–9
Torr. The XPS data were analyzed using the Hawk Data Analysis software
(V7.04.04) and CasaXPS software.

47

Focused ion beam (FIB) preparation

The black film was deposited on highly oriented pyrolytic graphite (HOPG) by immersing the
plate (1 cm

´ 1 cm, SPI-1 grade) face down into the reaction mixture. The material was washed
with water and ethanol, and the organic solvents were evaporated in vacuum.

Scanning electron microscopy (SEM) and Focused Ion Beam (FIB) milling were performed on
FEI NOVA 600. Images were acquired using secondary electron detection. The microscope’s high
tension was set to 10kV. The FIB samples were prepared using standard cross-section preparation
techniques.

The black film was washed with water and ethanol and deposited on SiN TEM windows purchased
from TEMWindows.com. The TEM data was acquired using a JEOL 2100F microscope. The
images were acquired using an Orius SC 200D camera. The microscope’s high tension was set to
200kV. The sample’s beam sensitivity required the use of frame capture from a 20 frame/second
video stream of the camera’s output. The output was captured using VirtualDub software. Frame-
by-frame analysis was used to extract diffraction patterns from the sample. Diffraction patterns
with hexagonal symmetry were seen on many parts of the film and a characteristic diffraction
pattern is shown in the text.

Energy dispersive X-ray spectroscopy (EDS) was acquired using an EDAX Apollo SDD-10
detector attached to a multi-beam JEOL JIB-4500 and a take off angle of 29.05°. The beam
accelerating voltage was 30 kV. Spectral maps were acquired from multiple regions of the film,
with typical dwell times of 200 s at each point. The SEI image was 514

´ 400 pixels and the
maps were 256

´ 200 pixels.

Powder X-ray diffraction (PXRD) was performed on a Rigaku Ultima IV X-Ray diffractometer
in reflectance parallel beam/parallel slit alignment geometry. The measurement employed Cu K
line focused radiation at 1760 W (40 kV, 44 mA) power and a Ge crystal detector fitted with a 0.6
mm radiation entrance slit. Samples were mounted on zero-background sample holders or highly
oriented pyrolytic graphite (HOPG) plates (5 cm

´ 1 cm

´ 0.3 cm) under a nitrogen atmosphere.
Samples were observed using a 0.04° 2 step scan from 2.0 – 60.0° with an exposure time of 0.4
s per step. No peaks could be resolved from the baseline for 2 > 35°.
48

2.4.4 Electrochemical methods

Electrochemistry experiments were carried out using a Pine potentiostat. Platinum wire used for
the electrochemical studies was purchased from Alfa Aesar. The electrochemical experiments
were carried out in three electrode electrochemical cells under nitrogen atmosphere using glassy
carbon (GC) or highly oriented pyrolytic graphite (HOPG) modified electrodes as the working
electrode. The diameters of the working electrodes are 3 and 9.3 mm for GC and HOPG,
respectively. A platinum wire, placed in a separate compartment, connected by a Vycor tip, and
filled with the electrolytic solution (0.1 M NaClO4 for aqueous media and 0.1 M nBu4NPF6 for
organic media) was used as auxiliary electrode. The reference electrode, placed in a separate
compartment and connected by a Vycor tip, was based either on an aqueous Ag/AgCl/saturated
KCl electrode (for aqueous media), or a Ag wire electrode (for organic media). The reference
electrode in aqueous media was calibrated externally relative to ferrocenecarboxylic acid (Fc-
COOH) at pH 7, with the Fe
3+/2+
couple at 0.28 V vs Ag/AgCl. All potentials reported in this paper
were converted to standard hydrogen electrode by adding a value of 0.205 V, or to reversible
hydrogen electrode (RHE) by adding a value of (0.205 + 0.059 pH) V. The reference electrode in
organic media was calibrated relative to ferrocene (Fc), with the Fe
3+/2+
couple at 0.38 V vs
Ag/AgCl.

Controlled-potential electrolysis measurements were conducted in a sealed two-chambered H cell
where the first chamber held the working and reference electrodes in 65 mL of 0.1 M NaClO4 (aq)
at the corresponding pH, and the second chamber held the auxiliary electrode in 25 mL of 0.1 M
NaClO4 (aq). The two chambers were separated by a fine porosity glass frit. The reference
electrode was placed in a separate compartment and connected by a Vycor tip. Glassy carbon plates
(5 cm

´ 1 cm

´ 0.3 cm; Tokai Carbon USA) were used as the working and auxiliary electrodes.
The reference electrode was a Ag/AgCl/saturated KCl(aq) electrode separated from the solution
by a Vycor frit.

The aqueous solutions used in the electrochemical experiments have been prepared as follows. For
the pH 0.8 and 1.0 solutions, NaClO4 (1.405 g) was dissolved in water (100 mL). 70% HClO4
solution (0.9 mL) was added to reach the desired pH. For the pH 1.3 solution, H2SO4 (0.534 mL,
c = 98%) was dissolved in water (200 mL) and the pH was measured. For the pH 2.6 solution,
49

citric acid (3.458 g) and Na2HPO4 (1.505 g) were dissolved in water (200 mL). For the pH 4.4
solution, NaOAc (1.605 g) was dissolved in water (200 mL). Acetic acid (1.2 mL) was added to
reach the desired pH. For the pH 7.1 solution, NaH2PO4 (0.468 g) and Na2HPO4 (1.637 g) were
dissolved in water (100 mL). For the pH 10.0 solution, NaHCO3 (0.339 g) and Na2CO3 (0.632 g)
were dissolved in water (100 mL). The pHs of the solutions were measured with a benchtop VWR
Symphony pH meter. All solutions were purged with nitrogen for 15 minutes.

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spectroscopy and Harry's intuition. Coordination Chemistry Reviews 2011, 255 (7), 837.
(40) Liu, S.-G.; Liu, Y.-Q.; Zhu, D.-B. Preparation and spectral properties of new molecular
conductors based on anion salt. Synthetic Metals 1997, 89 (3), 187.
(41) Zhou, S.; Ichimura, K.; Inokuchi, H. X-Ray photoelectron spectroscopy characteristics of
a novel organic semiconductor BTQBT and its derivatives. Journal of Materials Chemistry
1995, 5 (10), 1725.
(42) Solis, B. H.; Hammes-Schiffer, S. Computational Study of Anomalous Reduction
Potentials for Hydrogen Evolution Catalyzed by Cobalt Dithiolene Complexes. Journal of
the American Chemical Society 2012, 134 (37), 15253.
(43) Pullen, S.; Fei, H.; Orthaber, A.; Cohen, S. M.; Ott, S. Enhanced Photochemical Hydrogen
Production by a Molecular Diiron Catalyst Incorporated into a Metal–Organic Framework.
Journal of the American Chemical Society 2013, 135 (45), 16997.
52

(44) Harnisch, J. A.; Angelici, R. J. Gold and platinum benzenehexathiolate complexes as large
templates for the synthesis of 12-coordinate polyphosphine macrocycles. Inorganica
Chimica Acta 2000, 300-302, 273.
(45) Sakamoto, R.; Kambe, T.; Tsukada, S.; Takada, K.; Hoshiko, K.; Kitagawa, Y.; Okumura,
M.; Nishihara, H. π-Conjugated Trinuclear Group-9 Metalladithiolenes with a
Triphenylene Backbone. Inorganic Chemistry 2013, 52 (13), 7411.

CHAPTER 3

Metallic Conductivity in a Two-Dimensional Cobalt Dithiolene
Metal−Organic Framework

A portion of this chapter has appeared in print:

Clough, A. J; Skelton, J. M.; Downes, C. A.; de la Rosa, A. A.; Yoo, J. W.; Walsh, A.; Melot*,
B. C.; Marinescu*, S. C., “Metallic Conductivity in a Two-Dimensional Cobalt Dithiolene
Metal–Organic Framework” J. Am. Chem. Soc., 2017, 139, 10863–1086
54

3.1 Introduction
The development of inexpensive yet highly efficient catalysts for solar-to-fuel energy conversion
is vital for mitigating the adverse effects that hydrocarbon fuels have on the environment.
1-3

Immobilization of molecular catalytic units is an attractive strategy for bridging the gap between
homogeneous and heterogeneous electrocatalysis for solar-to-fuel applications.
4
This approach
retains the desirable properties of molecular systems, like well-defined chemistries and clear
design principles, while also taking advantage of the robust efficiency of heterogeneous catalysts.
4

Metal–organic frameworks, a rapidly expanding class of nanoporous coordination polymers,
5
have
attracted growing attention because they sit at the interface between molecules and extended
solids, offering a mixture of high surface area-to-volume ratio and site-isolation of catalytic units,
all of which are indicators to systems with enhanced activities. Yet, the biggest challenge these
materials face for designing new electrocatalysts is rooted in their ability to efficiently transport
charge between the metals and their coordinating ligands.
6-10

In this regard, 2D frameworks
11-15
have been shown to exhibit high charge-carrier mobility because
of in-plane charge delocalization and extended π-conjugation within the sheets.
16-22
Recent studies
have demonstrated that nickel or copper 2,3,6,7,10,11-hexaiminotriphenylene (M3(HITP)2, M =
Ni, Cu) frameworks exhibit thin film conductivities on the order of 40 S·cm–1, whereas copper
benzenehexathiolate shows conductivities up to 1580 S·cm–1, which is the highest conductivity
reported for a coordination polymer.
23-27
Yet, the transport properties are usually the result of
inadvertent doping of an intrinsic semiconductor that can be synthetically difficult to control rather
than a truly metallic charge delocalization.
In a previous report, we showed that 2,3,6,7,10,11-triphenylenehexathiolate can be used to produce
a periodic 2D network of cobalt dithiolene units as illustrated in Scheme 3.1.
28
This phase displays
remarkable activity for the electrocatalytic H2 evolution from water,
28
and is one of the first
examples of an electrocatalytically active MOF.
29-33
This motivated us to investigate the source of
this high catalytic performance by characterizing its charge transport characteristics. Here, we
report the first observation of a temperature-dependent transition in a MOF from a semiconducting
to a metallic phase with decreasing temperature using a combination of thin film resistivity
measurements and density functional theory (DFT) calculations, with metallic conductivity
persisting to temperatures as high as 225 K.
55

Scheme 3.1. Structure of the 2D cobalt dithiolene framework [Co3[THT]2]
3-
, CoTHT (1),
studied here. Reprinted with permission from Ref. 71. Copyright 2017 American Chemical
Society.
3.2 Results and Discussion
CoTHT was prepared using a slightly modified procedure from previous reports.
28
The material
adopts a hexagonal 2D structure with long-range order in the ab plane as evidenced by the sharp
Bragg reflections in the synchrotron X-ray diffraction pattern shown in Figure 2. The pattern
exhibits prominent peaks at 1.2°, 2.4°, 3.2°, and 4.2°, which correspond to a significant degree of
coherence within the 2D sheets. The somewhat broader reflection at 7.2° corresponds to the [001]
reflection, suggesting less coherence between the sheets as is common for layered materials.
34,35

Regardless, the experimental diffraction pattern is a close approximation to simulations from the
Materials Studio suite of programs (Figure 3.1) using the idealized structure with layers that are
stacked in perfect registry (Figure 3.2). The optimized structure is best described using P6/mmm
as the space group with unit cell parameters of a = b = 22.52 Å and c = 3.3 Å (Figure 3.3).
56

Figure 3.1. Spacefilling illustration of the fragment used to generate the model unit cell. Reprinted
with permission from Ref. 71. Copyright 2017 American Chemical Society.

Figure 3.2. Model structures of the eclipsed pattern. Reprinted with permission from Ref. 71.
Copyright 2017 American Chemical Society.

57

Figure 3.3. Experimental and simulated PXRD patterns of CoTHT. Reprinted with permission
from Ref. 71. Copyright 2017 American Chemical Society.

Gas sorption isotherms performed on CoTHT reveal a Brunauer-Emmett-Teller (BET) surface
area of 370 m
2
g
–1
(Figure 3.4), which is similar to that of the previously reported platinum
analogue.
21
Temperature-dependent susceptibility measurements show a response characteristic of
localized moments. A fit of the high temperature magnetic susceptibility data (200K–300K) to the
Curie-Weiss equation (χ=C/(T-θCW)+ χ0) yields an effective paramagnetic moment, µeff = 1.55 µB
per formula unit, θCW = –34K, and χ0= 3.57×10
–4
emu mol
–1
Oe
–1
(Figure 3.5). The positive
temperature-independent term reflects a background paramagnetic signal that could result from
the presence of some charge delocalization or more simply from second-order Zeeman effects.
36

The moment is slightly reduced from what is expected for an S = 1/2 state, which should show a
theoretical moment of 1.87 µB. Such an underestimate is often ascribed to orbital quenching due
to covalency, and indicates the presence of only one unpaired spin per formula unit. Given the
presence of three square-planar cobalt ions per unit cell, this is consistent with two thirds of the
Co adopting the trivalent state (S = 0) with one exhibiting a formal divalent (S = 1/2) state, which
is significant since mixed oxidation states are often associated with charge delocalization.
37,38

58

Figure 3.4. Nitrogen sorption isotherms performed on CoTHT at 77 K reveal a Brunauer-Emmett-
Teller surface area of 370(5) m
2
g
–1
. Reprinted with permission from Ref. 71. Copyright 2017
American Chemical Society.

Figure 3.5. Magnetic studies performed on the cobalt dithiolene framework CoTHT. (a) Magnetic
susceptibility data for CoTHT; (b) Curie-Weiss fit to the data; and (c) Magnetization (M) versus
applied magnetic field (H) data for CoTHT. Reprinted with permission from Ref. 71. Copyright
2017 American Chemical Society.

The temperature-dependent resistivity of CoTHT was measured using a four-point Van der Pauw
geometry on a pressed pellet of CoTHT with a thickness of 0.24(2) mm. Graphitic carbon paint
(Alfa Aesar) was used to create Ohmic contacts, as verified by the linear I-V trace shown in Figure
3.6. InGa eutectic and silver paint were also used to create Ohmic contacts and gave qualitatively
similar results; however, the carbon paint was the best at maintaining physical contact over the
entire temperature range of interest. At 300 K, the bulk resistivity was determined to be 0.720(7)
k-cm, corresponding to a conductivity of 1.4  10
–3
S·cm
–1
, which is in line with values reported
59

for the platinum analogue.
21
Given the highly anisotropic nature of the pressed pellet of CoTHT,
this relatively low value is likely associated with the random orientation of the powder and is likely
exaggerated by grain-boundary scattering between the sheet-like particles.

Figure 3.6. Room temperature I-V trace of solid CoTHT pressed in a pellet of 0.24(2) mm
thickness, which displays Ohmic response between –2.0 and +2.0 V. Reprinted with permission
from Ref. 71. Copyright 2017 American Chemical Society.

An exponential rise in the resistivity of the pellet is seen between 300 K and 170 K, as would be
expected for a semiconducting sample where transport is dominated by thermally populated
carriers that must overcome a hopping barrier for conduction (Figure 3.7a).
39
An Arrhenius fit to
the data suggests an activation energy on the order of 173 meV (Figure 3.8). More interestingly, a
decrease in the resistivity is subsequently seen between 130 K and 50 K, suggesting a transition to
a metallic phase where scattering of the carriers is dominated by lattice vibrations. The transition
is fully reversible, with no signs of hysteresis as illustrated in Figure 3.9, suggesting it is second
order in nature, unlike the metal-to-insulator transitions in materials like VO2 that are associated
with structural deformations.
40-42
To further confirm the absence of a coherent structural distortion,
variable temperature synchrotron powder X-ray diffraction studies were performed between 300
and 100 K. As indicated in Figure 3.10–3.12, the only significant change to the diffraction patterns
is a slight shift in the [001] reflection from 7.11° (300 K) to 7.15° (100 K), which corresponds to
a 0.02 Å contraction of the interlayer spacing or a roughly 1% change.
60

Figure 3.7. Variable-temperature resistivity data for (a) solid CoTHT pressed in a pellet of
0.24(2) mm thickness (yellow, scaled down 10
5
) and films of CoTHT with thicknesses of
0.10(1) (black), 0.12(1) (red), and 0.20(2) (blue) m deposited on glass supports; or (b) films of
CoTHT with thicknesses of 0.5(1) m before (blue) and after (red) a two-hour exposure under
vacuum at 90 °C. Insert (a): an SEM image of film CoTHT. Reprinted with permission from
Ref. 71. Copyright 2017 American Chemical Society.

Figure 3.8. Arrhenius plot of the high temperature conductivity data of solid CoTHT pressed in a
pellet of 0.24(2) mm thickness, displaying an activation energy (Ea) of 173 meV. Reprinted with
permission from Ref. 71. Copyright 2017 American Chemical Society.
61

Figure 3.9. Overlay of the variable-temperature resistivity data for solid CoTHT pressed in a
pellet of 0.24(2) mm upon cooling (blue) and warming (red). The resistivity data of the pressed-
pellet was measured at temperatures ranging from 300 to 60 K. Below 60 K the carbon paint
stopped maintaining physical contact with the sample. Reprinted with permission from Ref. 71.
Copyright 2017 American Chemical Society.

Figure 3.10. Synchrotron variable temperature powder X-ray diffraction (PXRD) patterns of solid
CoTHT collected at 100 (blue), 200 (green), 280 (orange), and 295 (red) K. No peaks could be
resolved from the baseline for 2 > 12°. Reprinted with permission from Ref. 71. Copyright 2017
American Chemical Society.

62

Figure 3.11. Overlay of the variable temperature PXRD patterns of CoTHT focusing on the [100]
reflections collected at 100 (blue), 200 (green), 280 (orange), and 295 (red) K. Reprinted with
permission from Ref. 71. Copyright 2017 American Chemical Society.

Figure 3.12. Overlay of the variable temperature PXRD patterns of CoTHT focusing on the [001]
reflections collected at 100 (blue), 200 (green), 280 (orange), and 295 (red) K. Reprinted with
permission from Ref. 71. Copyright 2017 American Chemical Society.

To explore this transition further, films of CoTHT were deposited on glass supports with scanning
electron microscopy (SEM) images indicating smooth surfaces and good coverage (Figures 3.7a
insert and 3.13). The thicknesses of the films were determined using atomic force microscopy
(AFM) and ranged from 0.1 and 0.2 m (Figures 3.14–3.15) with the resistivity being measured
in an identical fashion to the pellets described earlier (Figures 3.7 and 3.16-3.18). For films with a
thickness of 0.20 m, the resistivity at 300 K was 31 -cm, corresponding to a conductivity of 3.2
 10
–2
S·cm
–1
, which is a full order of magnitude higher than the conductivity of the pellet as
expected when grain boundaries between the sheet-like particles are reduced. The temperature-
dependent resistivity data of the films show a similar semiconductor-to-metal transition (Figure
3.7a). Interestingly, the transition temperature exhibits a strong correlation with film thickness,
63

with thinner films displaying a higher temperature transition to a metallic state and thicker films
remaining semiconducting to lower temperatures (Table 3.1).

Figure 3.13. Scanning Electron Microscope (SEM) images of the cobalt dithiolene film CoTHT
on glass support (15 kV accelerating voltage). Reprinted with permission from Ref. 71. Copyright
2017 American Chemical Society.

Figure 3.14. Atomic Force Microscopy (AFM) studies of the cobalt dithiolene film CoTHT of
0.20(2) m thickness on glass. Reprinted with permission from Ref. 71. Copyright 2017 American
Chemical Society.

Figure 3.15. Atomic Force Microscopy (AFM) studies of the cobalt dithiolene film CoTHT of
0.12(1) m thickness on glass. Reprinted with permission from Ref. 71. Copyright 2017 American
Chemical Society.

64

Figure 3.16. Images of the samples – pellet (a) or films of CoTHT (b) and (c) – and their puck
assemblies for conductivity measurements. Graphitic carbon paint was used to create Ohmic
contacts. Reprinted with permission from Ref. 71. Copyright 2017 American Chemical Society.

Figure 3.17. Typical variable temperature I-V traces of the cobalt dithiolene film CoTHT on a
glass support, displaying Ohmic response between –0.1 and +0.1 V. Reprinted with permission
from Ref. 71. Copyright 2017 American Chemical Society.

Figure 3.18. Variable-temperature resistivity data for films CoTHT with thicknesses of 0.10(1)
(black), 0.12(1) (red), and 0.20(2) (blue) m deposited on glass supports. The resistivity data of
the film with a 0.10(1) m thickness was measured at temperatures ranging from 350 to 180 K.
Below 180 K the carbon paint stopped maintaining physical contact with the sample. Reprinted
with permission from Ref. 71. Copyright 2017 American Chemical Society.
65

Figure 3.19. Overlay of the variable-temperature resistivity data for the cobalt dithiolene film of
CoTHT with a thickness of 0.20(2) m deposited on glass supports, upon cooling (blue) and
warming (red). Reprinted with permission from Ref. 71. Copyright 2017 American Chemical
Society.

Figure 3.20. Arrhenius plots of the high temperature conductivity data of the cobalt dithiolene
film CoTHT of 0.10(1) (red) and 0.20(2) (blue) m thicknesses on glass supports, displaying
activation energies (Ea) of 159 and 118 meV, respectively. Reprinted with permission from Ref.
71. Copyright 2017 American Chemical Society.

66

Table 3.1. Selected parameters for the variable-temperature resistivity analyses performed on a
pressed pellet or films of CoTHT. Tmetallic represents the lowest temperature at which the
transition between a semiconducting to metallic conductivity occurs. Reprinted with permission
from Ref. 71. Copyright 2017 American Chemical Society.

The nanoporous nature of MOFs is well-known to result in a significant amount of solvent being
trapped within the channels.
5
To investigate if the presence of trapped solvent has an influence on
the transition temperature, the variable temperature resistivity data of a film of 0.5(1) m thickness
was measured before and after a two-hour exposure under vacuum at 90 °C (Figures 3.7b, 3.21–
3.22). These measurements clearly indicate that the films with less solvent display an increase in
the metallic transition temperature from 105 to 130 K. XPS studies before and after the
conductivity experiments show no significant changes suggesting that the films are not
significantly altered by the thermal treatment (Figure 3.23).

Figure 3.21. Overlay of the variable-temperature resistivity data for film CoTHT with a thickness
of 0.5(1) µm before (a) and after (b) a two-hour exposure under vacuum at 90 °C, upon cooling
(blue) and warming (red). Reprinted with permission from Ref. 71. Copyright 2017 American
Chemical Society.

Entry Sample Thickness
Tmetallic
(K)
Resistivity (k-
cm) at 300 K
Conductivity
(mS·cm
–1
) at 300 K
Ea
(meV)
1 Pellet 0.24(2) mm 135 0.720 1.4 173
2 Film 0.10(1) m 225 0.048 21 159
3 Film 0.12(1) m 222 0.167 6.0 –
4 Film 0.20(2) m 180 0.031 32 118
67

Figure 3.22. Arrhenius plot of the high temperature conductivity data of film CoTHT with a
thickness of 0.5(1) µm before (blue) and after (red) a two-hour exposure under vacuum at 90 °C,
displaying activation energies (Ea) of 81 (before) and 114 (after) meV. Reprinted with permission
from Ref. 71. Copyright 2017 American Chemical Society.

68

Figure 3.23. X-ray photoelectron spectroscopy (XPS) data of the cobalt dithiolene film CoTHT
collected at room temperature (1) before and (2) after conductivity measurements. XPS data from
top: (a) Co 2p, (b) S 2s, and (c) S 2p. Reprinted with permission from Ref. 71. Copyright 2017
American Chemical Society.

On closer inspection of the transition, there are clearly two closely spaced maxima in the resistivity
data. While unusual, this type of transition has been observed in glassy charge transfer salt κ-
(BEDT-TTF)2Cu[N(CN)2]Br (where BEDT-TTF = bis-ethylenedithiotetrathiafulvalene).
43,44
In
this instance, the authors ascribed the unusual temperature dependence to a strong contribution
from lattice vibrations at higher temperatures in combination with highly anisotropic changes to
their in-plane lattice parameters. Both of these effects are likely at play in CoTHT as the
temperature-dependent synchrotron diffraction data shows a substantially more pronounced
69

change to the inter-layer spacing compared to the distances within the sheets. Considering that the
transitions in the resistivity are not reflected in the magnetic susceptibility, they are unlikely to be
associated with an in-plane structural distortion, which should alter the coupling between the spins.
Furthermore, the removal of solvent from the material would most likely result in a tighter packing
of crystallites and more interfacial contact between the sheet-like particles. Thus, our experimental
data suggests that changes in the vibrational modes, interlayer spacing and morphological changes
(such as contact at the grain boundaries) that interfere with inter-sheet interactions have the most
significant impact on the conductivity of the material.

In order to rationalize these observations in terms of the electronic structure of the framework,
DFT calculations were performed (see computational modeling section). The calculated band
dispersion and density of states (DOS; Figure 3.24) indicates the compound is actually a semi-
metal, with a small DOS at the Fermi energy (EF) in both spin channels. There is a large dispersion
of ~2 eV along the Γ-A line in the Brillouin zone, corresponding to the c-axis in real space. The
smallest calculated carrier effective mass for the metallic bands of 0.29 me (see computational
modeling section) suggest facile transport along this direction, and contrasts with a minimum
effective mass of 1.27 me along the in-plane directions. The bands making up the metallic states
correspond to π-type crystal orbitals centered on the metal ions and ligand S atoms, which explains
the large dispersion and suggests that the primary mechanism for conductivity is through
conductive pathways along the c-axis. The less facile in-plane transport is also supported by a
Bader-charge analysis (see computational modeling section).
45

70

Figure 3.24. Calculated electronic dispersion and density-of-states curve for CoTHT. Reprinted
with permission from Ref. 71. Copyright 2017 American Chemical Society.

To further investigate the stacking mode in CoTHT, potential energy surface (PES) calculations
were performed on a bilayer model with offsets of up to 4 Å along the a and b axes (Figure 3.25).
These studies show that the fully-eclipsed AA structure is the most energetically favored, although
a relatively shallow local minimum is present at offsets of ~1.75 Å along one or both axes.
Displacements of ±0.25 Å along either or both axes would easily be possible given the thermal
energy available at 300 K. These findings also offer some mechanistic insight into the temperature-
induced semiconductor-to-metal transition. Thermal expansion along the c-direction, or stacking
faults leading to misalignment of the layers, were found to introduce a gap in the conduction states
(Figures 3.31–3.34), which, in combination with a change in the Fermi level, or other factors such
as the behavior of the guest molecules in the pore, could play a role in the transition to a narrow-
gap semiconductor.

71

Figure 3.25. Contour map of the potential energy surface for offsets of alternate layers along the
a and b axes in CoTHT. Reprinted with permission from Ref. 71. Copyright 2017 American
Chemical Society.

Metallic conductivity has been suggested previously in a nickel benzenehexathiolate framework
using first-principles band structure calculations; however, conductivity measurements on a single
microflake revealed semiconducting behavior with a small activation energy (Ea) of 26 meV.
25

The discrepancy was attributed to structural disorder in the sample.
25
Moreover, DFT calculations
performed on Ni3(HITP)2 framework suggested that the bulk form is metallic, whereas the
monolayer form showed a small band gap of 0.25 eV.
27
The most energetically favored structure
for Ni3(HITP)2 was reported to be an AB slipped-parallel stacking mode wherein one layer was
slipped relative to the neighboring one by 1.8 Å along the a or b vectors.
23
Additionally, metal
substitution was shown by DFT studies to promote or change the electronic properties of these 2D
frameworks from semiconducting to metallic.
27,45-48

3.3 Conclusions

In summary, we have investigated the temperature-dependent resistivity of a cobalt 2,3,6,7,10,11-
triphenylenehexathiolate framework. Variable temperature resistivity studies performed on a
pressed-powder pellet indicate a semiconducting phase between 300 K and 170 K, followed by a
transition to metallic behavior at temperatures below 130 K, which has been unprecedented in
MOFs. A similar transition is observed for films, with the transition temperature being highly
72

dependent on the film thickness. Electronic-structure calculations support the experimentally
observed complex metallic conductivity, with the highest mobility pathways occurring between
the sheets. The temperature-dependence of the resistivity exhibits multiple maxima, which
suggests that contributions from stacking faults, local molecular vibrations, and the behavior of
solvent molecules in the pores may all be convoluted together to produce a complex mechanism
for scattering the charge carriers.
Overall, these results identify the first experimentally observed MOF that exhibits band-like
metallic conductivity, and highlights the importance of external factors like guest molecules and
film morphology in obtaining highly conductive 2D frameworks. We expect the design principles
discovered in these fundamental studies to have a profound impact in understanding the charge
transport characteristics of MOFs, leading to new materials with impressive electrical properties.

3.4 Experimental Details

3.4.1 General Considerations

All manipulations of air and moisture sensitive materials were conducted under a nitrogen
atmosphere in a Vacuum Atmospheres drybox or on a dual manifold Schlenk line. The glassware
was oven-dried prior to use. Acetonitrile and dichloromethane were degassed with nitrogen and
passed through activated alumina columns and stored over 4 Å Linde-type molecular sieves. Ethyl
acetate, water, and ethanol were placed under vacuum and refilled with nitrogen (10 ). Deuterated
solvents were dried over 4 Å Linde-type molecular sieves prior to use. Elemental analyses were
performed by Robertson Microlit Laboratories, 1705 U.S. Highway 46, Suite 1D, Ledgewood,
New Jersey, 07852. All the chemical regents were purchased from commercial vendors and used
without further purification. The ligand 2,3,6,7,10,11-triphenylene-hexathiol (THT) was prepared
according to the reported procedures.
49
Water was deionized with the Millipore Milli-Q Synergy
system (18.2 M·cm resistivity). All other chemical regents were purchased from commercial
vendors and used without further purification.

3.4.2 Synthesis of CoTHT

The cobalt dithiolene framework CoTHT was prepared according to the reported procedures.
28
A
400 mL glass jar was charged with an aqueous solution of CoCl2·6H2O (60 mg, 0.25 mmol, 2.5
73

mM, 100 mL volume). Separately, a suspension of 2,3,6,7,10,11-triphenylene-hexathiol (THT)
(10 mg, 0.024 mmol) in N-Methyl-2-pyrrolidone (NMP) (0.2 mL) was then diluted with ethyl
acetate until the total volume of the suspension reached 20 mL, sealed, and briefly sonicated to
form an uniform suspension. Ethyl acetate (80 mL) was gently added to the aqueous solution to
create a liquid-liquid interface; the suspension of THT was then gently added to the ethyl acetate
layer and the jar was sealed and allowed to stand overnight. A black film appeared at the liquid-
liquid interface over 5 days. The film was deposited on glass supports by pulling the substrate up
through the film. The deposited films were then washed with water and allowed to evaporate to
dryness. Alternatively, the black solid CoTHT was collected by filtration and washed with water
and methanol for bulk powder analyses.

X-Ray Photoelectron Spectroscopy (XPS)

XPS data were collected using a Kratos AXIS Ultra instrument. The monochromatic X-ray source
was the Al K α line at 1486.7 eV, and the hybrid lens and slot mode were used. Low-resolution
survey spectra were acquired between binding energies of 1–1200 eV. Higher-resolution detailed
scans, with a resolution of 0.1 eV, were collected on individual XPS regions of interest. The sample
chamber was maintained at < 9  10
–9
Torr. The XPS data were analyzed using the CasaXPS
software.

Powder X-ray diffraction (PXRD) studies were performed on a Rigaku Ultima IV X-Ray
diffractometer in reflectance parallel beam/parallel slit alignment geometry. The measurement
employed Cu K line focused radiation at 1760 W (40 kV, 44 mA) power and a Ge crystal detector
fitted with a 2 mm radiation entrance slit. Samples were mounted on zero-background sample
holders and were observed using a 0.08° 2 step scan from 2.0 – 40.0° with an exposure time of
0.4 s per step. No peaks could be resolved from the baseline for 2 > 35°.

High resolution synchrotron powder X-ray diffraction data was collected using the 11-BM
beamline mail-in program at the Advanced Photon Source (APS), Argonne National Laboratory,
with an average wavelength of 0.414575 Å. Discrete detectors covering an angular range from 0.5
to 30º 2θ are scanned over a 34º 2θ range, with data points collected every 0.001º 2θ and scan
74

speed of 0.01º/s. An Oxford Cryosystems Cryostream Plus device allowed for sample temperatures
to be controlled over a range of 100-295 K.

Modeling

Molecular modeling of CoTHT was carried out using the Materials Studio (version 8.0) suite of
programs by Accelrys. The molecular fragment used to generate the model is shown in Figure 3.1.
The unit cell was constructed starting with a primitive hexagonal unit cell with space group
P6/mmm using cell parameters a = b = 22.52 Å and c = 3.3 Å. The structure was optimized with
Materials Studio Forcite calculations using geometry optimization and universal forcefield
methods. The MS Reflex module was used to calculate the expected PXRD patterns. Line
broadening for crystallite size was not calculated. Comparison of the simulated and experimental
PXRD patterns verified the simulated structure.

Conductivity Measurements

Conductivity measurements were performed using a custom set up integrated into a 14T Quantum
Design Dynacool Physical Properties Measurement System. A Keithley 6220 Precision Current
Source (excitation currents of 1-50 nA) was used to trigger and control a Keithley 2182A
nanovoltmeter. In order to minimize errors associated with contact resistance and drift voltages, a
Keithley 2172 matrix switch equipped with a Keithley 6536 Hall effect card was used to alternate
the direction of the applied current. Because of difficulty associated with preparing samples with
uniform dimensions, all measurements were performed in a four-point probe Van der Pauw
geometry. Copper wire contacts were attached to the films using a conductive carbon paint and
soldered onto a Quantum Design puck with resistivity option. All measurements were performed
under a reduced pressure of ~10 torr.

Scanning Electron Microscopy (SEM)

SEM images were collected using a JEOL-7001F operating at 15 kV with 5 nA of probe current.

Atomic Force Microscopy (AFM)

AFM topography images were collected in tapping mode using an Agilent 5420 SPM instrument
operating in tapping mode. The probe tips were Tap300-G Silicon AFM probes (resonant
75

frequency 300 kHz, force constant 40 N/m) purchased from Budgetsensors.com and aligned prior
to use. Images were collected with a scan rate of 0.1 lines per second and over an area of 40 µm.
All samples were imaged under one atmosphere of air at room temperature. We are grateful to Dr.
John Chen and Prof. Mark Thompson for assistance with the AFM measurements.

Gas sorption analysis

Brunauer-Emmett-Teller (BET) measurements were performed on a Nova 2200e surface area and
pore size analyzer (Quantachrome Instruments, Inc.). Samples were degassed for 2 h at 150 °C in
vacuo prior to measurements. We are grateful to Dr. Sean Culver and Prof. Richard Brutchey for
assistance with the BET measurements.

3.4.3 Computational Modeling

Density-functional theory calculations on the crystal structure of CoTHT were carried out within
the plane-wave pseudopotential formalism, as implemented in the Vienna ab initio Simulation
Package (VASP) code.
50

We employed the PBEsol exchange-correlation functional
51
with the DFT-D3 dispersion
correction.
52
The ions were modelled with projector augmented-wave (PAW) pseudopotentials
53,54

treating the H 1s, C 2s and 2p, S 3s and 3p and Co 4s, 3d and 3p electrons as valence states. A
kinetic-energy cutoff of 800 eV was used to define the plane-wave basis, and the electronic
Brillouin zone was integrated using Γ-centered Monkhorst-Pack k-point meshes
55
with 1×1×5 and
2×2×15 subdivisions for geometry optimisations and electronic-structure calculations,
respectively. Careful convergence testing found that the 1×1×5 mesh was sufficient to converge
the absolute value of the total energy and unit-cell pressure to within 1 meV per atom and 1 kbar
(0.1 GPa), respectively, while the denser mesh was required to obtain a high-quality electronic
density of states and an accurate Fermi energy.

Tolerances of 10
–8
eV and 10
–2
eV Å
–1
were applied during the optimisation of the Kohn-Sham
wavefunctions and atom positions and cell parameters, respectively. The PAW projection was
performed in real space. The precision of the charge-density grids was set automatically to avoid
aliasing errors. For the electronic-structure calculations, a Gaussian smearing with a width, 𝜎 , of
76

0.05 eV was used to broaden the density of states (DoS), and band dispersions were obtained by
calculating eigenvalues non-self consistently at strings of k-points along the high-symmetry paths
through the Brillouin zone.

Magnetic Structure and Equilibrium Geometry

Given the complex magnetic structure of CoTHT, we performed a series of calculations with
different initial magnetic moments and compared the spin density and total energies of the
resulting electronic groundstates.

We performed two initial calculations starting with 0 and 1 unpaired spin(s) on each ion (i.e. on
the Co ions and the ligand atoms), and allowed the spin density to relax during the wavefunction
optimisation. We then performed a series of additional calculations starting from one of seven
initial moments corresponding to Co
2+
and Co
3+
ions with unpaired spin density distributed among
the bonded S atoms as appropriate for a thiolate anion/radical configuration (Figure 3.26). In these
calculations, the total magnetic moment was either fixed to the initial value or allowed to relax
during the minimisation.

We found five unique magnetic configurations, of which the lowest energy had a total moment of
3.8 BM per unit cell with approx. 0.9 BM (one unpaired electron) associated with each Co ion and
the remaining density associated with the S and C atoms of the ligands. A larger proportion of the
unpaired spin density on the ligand was assigned to S than to C, and each S atom had an equal
magnetic moment. This state can thus be assigned as Co
2+
ions with each of the S atoms on the
ligand being of mixed ion/radical character.

77

Figure 3.26. Initial local magnetic moments and nominal total moments of seven magnetic
configurations tested during preliminary calculations to determine the optimum magnetic
configuration of the unit cell of CoTHT. Reprinted with permission from Ref. 71. Copyright 2017
American Chemical Society.

We next performed further calculations in which the seven local magnetic configurations in Figure
3.26 were arranged in a frustrated antiferromagnetic configuration, with the sign of the initial
moments reversed on one of the three Co ions in the unit cell. These yielded a further four unique
magnetic configurations, none of which were lower in energy than the ferromagnetic state.

Optimisation of the atomic positions and unit-cell parameters with the most energetically-
favourable initial configuration yielded a structure close to the experimental one (𝑎 opt
= 23.133
Å, 𝑐 opt
= 3.140 Å; c.f. 𝑎 = 22.52 Å, 𝑐 = 3.3 Å from the experimental measurements). Moreover,
the spin density relaxed during the geometry optimisation to a final value of 1.8 BM, which is in
good agreement with the moment of 1.5 BM per formula unit measured experimentally. Analysis
of the spin density yielded local moments of ~0.6 BM on each Co ion, with negligible spin density
on the ligands (Figure 3.27).

A single-point calculation on the optimised cell starting with an antiferromagnetic arrangement of
the Co spins produced a configuration 72 meV per Co ion higher in energy than the ferromagnetic
arrangement, confirming the latter to be the groundstate.

78

Finally, we also performed calculations on a model with a –3 𝑒 charge per unit cell on the
framework and a homogenous compensating background charge, which would allow for three
Co
3+
ions with dianionic dithiolate ligands. This configuration proved to be unstable during
geometry optimisation, however, and resulted in layer separation (𝑐 opt
= 14.963 Å). This could be
indicative of charges on the framework being unstable, but may also reflect the homogenous
background charge being a poor model for explicit charge-compensating cations in the pores.
We therefore opted to perform calculations on the uncharged model with the ferromagnetic spin
configuration.

Figure 3.27. Calculated spin density of the optimised model of CoTHT. This image was prepared
using the VESTA software.
56
Reprinted with permission from Ref. 71. Copyright 2017 American
Chemical Society.

Electronic Structure
79

As described in the text, electronic-structure calculations on CoTHT revealed it to be a semi metal,
with a small density of states at the Fermi energy and a large dispersion along the real-space 𝑐
direction.

We estimated the carrier effective masses 𝑚 ∗
from the curvature of the band dispersion according
to:
1
𝑚 ∗
=
1
ℏ
2
𝜕 2
𝐸 (𝑘 )
𝜕 𝑘 2

(1)

where ℏ is the reduced Planck constant. The bands contributing to the metallic states were
identified as those crossing the Fermi energy along the Γ-𝐴 line in the electronic Brillouin zone
within 25 meV of the Fermi energy. The energies of those bands along the 𝑀 -Γ, Γ-𝐴 and 𝐴 -𝐻
segments were then fit to quadratic polynomials of the form 𝐸 (𝑘 ) = 𝑎 𝑘 2
+ 𝑏𝑘 + 𝑐 , from which
the second derivative term in Eq. (1) can be extracted as 𝜕 2
𝐸 (𝑘 ) 𝜕 𝑘 2
⁄ = 2𝑎 .

Whereas the bands along the 𝑀 -Γ and 𝐴 -𝐻 segments are well fit by single quadratic functions, the
more complex dispersion along the Γ-𝐴 segment is highly non-parabolic. We therefore fitted this
segment in a “piecewise” fashion to multiple functions, and extracted “local” second derivatives
for segments close to the Fermi energy. Some bands along this segment also exhibit sharp band
crossings that cannot be fitted to a quadratic function; to avoid including these in the effective-
mass calculations, we omitted fits where the maximum absolute error of the fit was greater than
25 meV. The fits used to calculate the effective masses are overlaid on the band structure from
Figure 3.25 in the text in Figure 3.28.

80

Figure 3.28. Calculated electronic dispersion and density-of-states curve for CoTHT. The blue
and red lines show the spin-up and spin-down components of the electronic structure, respectively.
The thick black overlaid lines illustrate the fitting of the bands contributing to metallic states to
quadratic polynomials in order to estimate the carrier effective masses, as described in the text.
Reprinted with permission from Ref. 71. Copyright 2017 American Chemical Society.

From this analysis, the effective masses of the bands near the Fermi energy along the 𝑀 -Γ segment
range from 1.27 to 18.7 𝑚 𝑒 in the spin-up channel, while all of the bands in the spin-down channel
lie outside of the 25 meV threshold. Along the Γ-𝐴 segment, the calculated effective masses range
from –0.42 to –1.52 𝑚 𝑒 and 1.43 to 8.04 𝑚 𝑒 in the spin-up channel, and -0.43 to –0.59 𝑚 𝑒 and
0.29 to 7.01 𝑚 𝑒 in the spin-down channel. None of the bands along the 𝐴 -𝐻 segment fall within
25 meV of the Fermi energy.

To establish the nature of the conductive states, we identified the partially-occupied bands and
calculated an orbital density plot showing the associated crystal orbitals (Figure 3.29). This
analysis shows that the valence orbitals consist mainly of Co d and S p orbitals, with a small
contribution from the ligand C p orbitals. This suggests that the primary conduction mechanism is
via channels formed of the metal and S valence orbitals.

Although the orbital density plots do not show the phasing of the orbitals, it can be inferred that
they are of 𝜋 symmetry, since the band energies reduce from Γ (in-phase overlap between unit
cells) to 𝐴 (antiphase overlap between unit cells along the c direction). The large dispersion of
these bands can therefore be explained by a strong interaction between metal centres in alternate
layers.
81

Figure 3.29. Orbital-density plot showing the crystal orbitals associated with the partially-
occupied bands in the optimised model of CoTHT. This image was prepared using the VESTA
software.
56
Reprinted with permission from Ref. 71. Copyright 2017 American Chemical Society.

Following the literature reports,
45
we also performed a Bader-charge analysis to estimate the
charge on the Co ions. We obtained a Bader charge of 0.51 𝑒 per atom, which, as per the literature
analysis, suggests a predominantly ionic character to the Co-S bonds, which is in keeping with
more facile transport along the c axis compared to the in-plane a and b directions.

Finally, although performing calculations with explicit charge-compensating cations in the pores
is infeasible given the lack of crystallographic order, to test the possible effects of reducing the
framework on the electronic structure, we recalculated the dispersion and DoS curves in Figure
3.25 in the text/Figure 3.28 with a charge of –3 𝑒 per unit cell and a compensating homogenous
background charge (Figure 3.30).
82

Figure 3.30. Calculated electronic dispersion and density-of-states curve for CoTHT with a
charge of –3 𝑒 per unit cell. As in Figure 3.28, the blue and red lines show the spin-up and spin-
down components of the electronic structure, respectively, and the overlaid black lines illustrate
the fitting to estimate the carrier effective masses. Reprinted with permission from Ref. 71.
Copyright 2017 American Chemical Society.

Comparison with Figure 3.25 in the text suggests that the negative charge leads to additional bands
around the Fermi level but, crucially, does not change the metallic nature, nor the large band
dispersion along the Γ-𝐴 segment. The calculated carrier effective masses are –0.59 and 1.37 𝑚 𝑒
in the spin-up channel and 0.75 to 6.71 𝑚 𝑒 in the spin-down channel, which are comparable to the
values obtained for the neutral cell.

Structural Flexibility

To explore the origin of the metal-to-semiconductor transition in CoTHT, we performed
additional calculations to study the effect of structural deformation on the electronic structure.

To investigate the effect of expansion along the c direction, we performed a series of single-point
energy calculations on the optimised model in which the c axis was lengthened by up to 15 Å. To
study the effect of layer misalignment, we prepared a 1×1×2 supercell expansion and performed
a series of single-point calculations in which one of the two layers was displaced with respect to
the other by up to 4 Å along the crystallographic a and b directions. To allow for changes to the
equilibrium interlayer spacing, at each a/b displacement we performed a series of calculations with
the length of the c axis adjusted by –0.3 to 0.5 Å in 0.1 Å steps. The resulting energy as a function
83

of c-axis length was then interpolated with a cubic spline, and the optimum interlayer spacing and
associated total energy obtained by locating the minimum with a binary search.
In the calculations on the single-layer models, we tried both fixing the magnetic moment to the
value obtained for the optimised structure (1.8 BM) and allowing it to optimise from an initial
configuration with one unpaired spin on each Co ion.

In the single-point calculations on the bilayer model, the k-point mesh was reduced to 1×1×3
subdivisions for the longer c axis. We found in preliminary calculations on a subset of the
structures with constrained magnetic moments that misaligning alternate layers had a substantial
influence on the spin distribution. Since to enumerate possible alternative magnetic configurations
at each data point as we did in our preliminary calculations was infeasible, we opted to allow the
magnetic moment to optimise freely in these calculations. For the same reason, we did not relax
the atomic positions, although since we do not change the lengths of the a and b axes in these
calculations, and we would expect the in-plane bonding to be quite rigid, we expect this to be a
reasonable approximation.

Figure 3.31. Energy change ∆𝐸 as a function of the interlayer spacing (c-axis length) of CoTHT.
Calculations were performed with the magnetic moment fixed to the lowest-energy value
determined for the equilibrium structure (𝑀 = 1.8 BM per unit cell; blue tringles) and with the
magnetisation allowed to freely optimise during the wavefunction minimisation (𝑀 = Free, red
circles). Reprinted with permission from Ref. 71. Copyright 2017 American Chemical Society.

The change in energy as a function of the interlayer spacing is shown in Figure 3.31. Allowing the
magnetisation to optimise freely leads to slightly lower energies as the layer separation increases,
84

with a maximum reduction of 94 meV per Co ion (7.8 %). For interlayer spacing larger than ~9 Å
the change in energy plateaus, giving interlayer interaction strengths of 1.2 /1.1 eV (116/106 kJ
mol
-1
) per Co ion with a fixed and free magnetic moment, respectively. These values are
comparable to a strong hydrogen bond. Given modest amounts of thermal energy (𝑘 B
𝑇 ≈ 25 meV
at 300 K), we predict a c-axis expansion of approx. 0.06-0.08 Å (1.9-2.5 %) to be accessible.

The change in total energy as a function of the offset of alternate layers along the a and b axes
forms a two-dimensional potential-energy surface (Figure 5 in the text; reproduced here as Figure
3.32). Unlike the study on the similar material Ni3(HTTP)2,
23
in this case we find the “eclipsed”
configuration to be the energetic minimum. There appear to be local minima at layer slips of ~1.75
Å along either or both axes, albeit at ~500 meV (83 meV per Co ion) above the equilibrium
structure. That the eclipsed conformation is favoured is not unexpected, since the same orbital
interactions that give rise to the metallic conductivity would presumably favour this arrangement.
Our analysis suggests that a small degree of misalignment should be easily accessible
energetically, with sufficient thermal energy at 300 K to allow the system to explore offsets of up
to 0.25 Å along either or both the two axes.

Figure 3.32. Contour plot of the total energy of a bilayer of CoTHT as a function of layer offset
along the crystallographic a and b axes, with the c-axis length optimised at each data point. A grid
of calculated points, marked by dashed lines, was used to generate the smooth surface by
interpolation. Reprinted with permission from Ref. 71. Copyright 2017 American Chemical
Society.

85

Figure 3.33 shows the variation in the calculated electronic density of states (DoS) as the layer
spacing is increased by 0.08, 0.15, 0.25 0.5 and 1 Å from the optimised equilibrium value,
corresponding to energetic costs of ~26.9, 101, 236, 794 and 1608 meV per unit cell, respectively.
Increasing the layer spacing causes a reduction in the density of states around the Fermi energy,
with a gap opening up around 1 eV in the conduction band. A narrowing of the bandwidth is also
evident, which may correspond to a reduction in the band dispersion and hence an increase in the
carrier effective mass and reduced mobility.

Figure 3.34 shows the calculated DoS of bilayer models with layer offsets of up to 1.75 Å.
Displacements of 0.25 Å along the a/b axes independently and a displacement of 0.25 Å along
both axes produced a near-identical DoS, in keeping with the very similar energy evident in Figure
3.9, and since the lowest-energy offset appears to be the a = b line, we focussed on simultaneous
offsets along both axes. Layer misalignment appears to have a similar effect to increasing the layer
spacing, with a gap in the conduction band clearly evident for the 1 and 1.75 Å offsets. However,
the requisite structural changes carry a much smaller energy penalty than increasing the layer
separation, plus there is the possibility of local structural distortions becoming transiently
“trapped” in the high-energy minima evident in Figure 3.9

Taken together, these results suggest that both changes to the interlayer spacing and layer
misalignment may play a role in the semimetal-to-semiconductor transition in CoTHT, although
given the significant energies associated with the larger structural distortions we would suggest
that other factors, such as the behaviour of guest molecules in the pores, are likely also to play a
role. One possibility is that interaction with guest molecules may shift the Fermi level, which, in
combination with the opening of a gap in the conduction band, could lead to (dominant)
semiconducting behaviour.

86

Figure 3.33. Calculated electronic density-of-states curves of CoTHT with different layer
spacings relative to the equilibrium value 𝑐 0
. The spin-up and spin-down components of the DoS
are shown as blue and red filled curves, respectively. Reprinted with permission from Ref. 71.
Copyright 2017 American Chemical Society.

87

Figure 3.34. Calculated electronic density-of-states curves of a bilayer model of CoTHT with
several displacements of the layers along the a and b axis relative to each other. The spin-up and
spin-down components of the DoS are shown as blue and red filled curves, respectively. Reprinted
with permission from Ref. 71. Copyright 2017 American Chemical Society.

88

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(45) Shojaei, F.; Hahn, J. R.; Kang, H. S. Mechanical and Electronic Properties of π-Conjugated
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(46) Chen, S.; Dai, J.; Zeng, X. C. Metal–organic Kagome lattices M3(2,3,6,7,10,11-
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CHAPTER 4

Inducing Room Temperature Metallic Conductivity in a MOF via Oxidation

93

4.1 Introduction

Metal-organic frameworks (MOFs) are crystalline nanoporous materials composed of
metal ions or clusters linked by organic ligands.
1-4
The hybrid organic/inorganic nature of
MOFs allows for synthetic tunability, leading to MOFs with varying pore sizes and
chemical environments, and therefore, different physical and chemical properties.
1-5
These
properties have led to applications primarily in gas storage and separation
6-8
and catalysis,
9

which take advantage of the inherent porosity and high surface area of these materials.
10

The use of MOFs in technologies that require charge transport, such as electronics
11,12
and
electrocatalysis,
13
has lagged behind due to the generally poor electrical conductivity of
MOFs. The weak covalent overlap between the metal and ligand orbitals results in flat,
low-dispersion electronic band structures with localized charge carriers, preventing
significant charge transport through the framework. This leads to materials with insulating
or large band-gap semiconducting behavior.
5,14,15
Efforts to reduce the barriers to charge
transport have included the addition of guest species,
16-18
doping,
19-21
and variation of the
metal center and its oxidation state.
22-26
These modifications can encourage through-
space
27,28
or through-bond
29
electronic transport and have led to MOFs with improved
conductivities, with one example reporting tunable conductivity over six orders of
magnitude.
17

Recently, the development of MOFs with redox active linkers has led to a breakthrough in
the field of intrinsically conductive MOFs.
11,12,15,30-33
Several two- and three-dimensional
(2D/3D) frameworks with planar, π-conjugated, and redox-active linkers, like
semiquinones/cathecolates,
34-42
diimines,
43-47
and dithiolenes,
19,21,48-59
have been reported
and display high conductivity values. However, while computational studies often predict
these 2D MOFs to be metallic,
49,60-63
the majority of the frameworks reported display a
decrease in conductivity upon cooling as expected for a semiconducting material where
transport is dominated by thermally-populated carriers. Metals, in contrast, display an
increase in conductivity upon cooling, as thermal energy scatters otherwise mobile valence
electrons. We have previously reported the temperature-dependent electronic conductivity
of a 2D cobalt 2,3,6,7,10,11-tripheylenehexathiolate (THT) MOF.
51
While the measured
conductivity at 300 K was modest (3.2⨯10
–2
S/cm), a complex transition from
94

semiconducting to metallic behavior was observed upon cooling, which represents the first
direct experimental observation of a MOF exhibiting band-like metallic conductivity. The
metallic conductivity was corroborated by density-functional theory (DFT) calculations,
which predicted semimetallic behavior through electronic states arising from interlayer
overlap of metal d and ligand p orbitals. This complex behavior was attributed to thermal
expansion along the c-direction, or stacking faults leading to misalignment of the layers,
which were found to introduce a gap in the conduction states.

Recently, an analogous iron THT MOF with ammonium charge-compensating cations,
(Fe3(THT)2(NH4)3), was studied using high-frequency terahertz photoconductivity and Hall
effect measurements. These studies indicate that Fe3(THT)2(NH4)3 displays
semiconducting behavior with record room-temperature carrier mobilities.
55
The calculated
band structure, which included explicit ammonium counterions, similarly predicts
semiconducting behavior with a bandgap of ~350 meV. Strong orbital hybridization is
observed between the d orbitals of Fe, the bis(dithiolene) moieties, and the triphenylene
units. While promising room-temperature mobilities were reported for this material, the
reported temperature-dependent resistivity studies indicate no evidence of band-like
metallic conductivity upon cooling, in contrast to the cobalt analogue.
51
This discrepancy
emphasizes the current knowledge gap in developing methods to control the nature of
charge transport (i.e., semiconducting vs. metallic conductivity) in MOFs. This
shortcoming is especially noteworthy in materials reported to date using
benzenehexathiolate and hexaiminobenzene linkers, which display decreases in their
conductivity upon cooling (semiconductor-like behavior), despite ultraviolet photoelectron
spectroscopy (UPS) studies revealing Fermi edges, which strongly suggest metallic
character.
45,49,58

Related studies of a copper benzenehexathiolate (CuBHT) MOF report a decrease in
conductivity upon cooling, consistent with metallic conductivity.
48
By using a modified
synthetic procedure, the crystallinity of the resulting CuBHT was improved as indicated by
powder X-ray diffraction (PXRD) studies.
54
Temperature-dependent resistivity studies of
the crystalline MOF indicate a decrease in resistivity upon cooling, as expected for a metal,
95

followed by the noteworthy observation of a superconducting transition at low
temperatures. These studies highlight the influence of crystallinity in dictating the transport
properties of 2D MOFs (semiconducting vs. metallic). Ideally, however, the transition from
semiconducting to metallic behavior would be controlled chemically and post-synthetically
for facile device preparation, rather than through a reliance on generating single-crystal
samples (a great challenge in the field of 2D MOFs). As a promising example of this goal,
the chemical reduction of an analogous silver benzenehexathiolate (AgBHT) 2D MOF was
shown by UPS studies to eliminate the observed Fermi edge of the pristine material,
suggesting a metallic-to-semiconducting transition, with a corresponding reduction
(~3,500-fold) in its electronic conductivity.
58
Despite these results, a mechanism by which
to control semiconducting-to-metallic transition temperature in 2D MOFs has not been
demonstrated.

Motivated by the knowledge that subtle chemical changes (metal identity, doping, guest
species, etc.) play key roles in determining the electronic properties of metal bis(dithiolene)
species and related 2D MOFs,
64-66
we study here a modified synthesis of the FeTHT
framework to generate a structural analogue to CoTHT. This modified protocol is informed
by DFT predictions and seeks to circumvent the inclusion of ammonium guest species and
generate materials with improved crystallinity as a means of inducing the semiconducting-
to-metallic transition observed in CoTHT. Temperature-dependent resistivity studies of the
as-prepared FeTHT framework confirm a transition from semiconducting to metallic upon
cooling, which is analogous to the behavior observed for the CoTHT system.
51
We further
demonstrate here the role of post-synthetic modification with a mild and abundant oxidant
(O2) in tuning the electronic properties of the FeTHT framework. Chemical oxidation of
the sample leads to an increase in the semiconducting-to-metallic transition temperature as
a function of O2-exposure, resulting in a material with metallic character at room
temperature. Through a combination of X-ray photoelectron spectroscopy (XPS) and
magnetic studies, we confirm that both the metal center and the ligand scaffold in FeTHT
are oxidized. These results demonstrate, for the first time, chemical control over the
semiconducting-to-metallic transition temperature in a MOF. The ability to bring about
drastic transformations in the electronic properties of these advanced materials presents a
96

major step towards enabling their utilization in broad applications ranging from
optoelectronics and chemiresistive sensing to electrocatalysis.

4.2 Results and Discussion

4.2.1 Synthesis and Chacterization

The 2D FeTHT MOF was synthesized using a liquid-liquid interfacial reaction as
previously reported for the cobalt analogue (Figure 4.1).
51,67
The trinucleating ligand
scaffold, triphenylene-2,3,6,7,10,11-hexathiol, was treated with N-Methyl-2-pyrrolidone
(NMP) in an ethyl acetate solution and sonicated. The generated suspension was gently
layered onto an aqueous solution of the iron(II) chloride, leading to the formation of a black
film of FeTHT at the liquid-liquid interface over the course of 5 days. The black film was
then deposited onto glass substrates as thin films or collected as a powder for bulk
measurements. This synthetic methodology is in contrast to that reported recently for
Fe3(THT)2(NH4)3, where a solution of Fe(acac)2 in chloroform was layered with an aqueous
solution of THT ligand with NH4OH as base.
55
The variations employed here serve to
modify the chemical formula unit by substituting the NH4
+
guest species with NMP, which
is subsequently removed by washing to generate the desired Fe3(THT)2(methanol)2
stoichiometry (as confirmed by elemental analysis). This approach provides a platform for
demonstrating the influence of guest species on bulk transport properties of the iron system,
as well as appropriately comparing the electronic behavior of the iron and cobalt systems.

The crystallinity of the FeTHT framework was confirmed by powder X-ray diffraction
(PXRD) using synchrotron radiation. FeTHT displays prominent peaks at 1.21°, 2.18°,
2.42°, 3.21°, and a broad peak at 7.09° (Figure 4.1). These peaks are similar to the ones
reported for the CoTHT framework, suggesting analogous structural environments for the
cobalt and iron systems.
51,67
The peak at 1.21° corresponds to the [100] reflection and is
indicative of a pore diameter of approximately 2.0 nm. This pore diameter is consistent
with the recently reported PXRD spectrum of Fe3(THT)2(NH4)3 (~1.9 nm), confirming that
the modified synthetic procedure employed here enables substitution of the guest species
without substantially altering the 2D structure of the material.
55
Figure 4.1 illustrates a
97

comparison of the experimentally observed diffraction pattern of FeTHT with the
simulated pattern of a model using the P6/mmm space group and unit cell parameters of a
= b = 22.52 Å and c = 3.34 Å (Figure 4.2). This diffraction data indicates good long-range
order in the ab plane and weak ordering along the c direction. The morphology of FeTHT
films was examined using scanning electron microscopy (SEM). Low magnification
images (Figure 4.1, inset) show few cracks, and higher magnification images (Figure 4.2)
reveal flat, sheet-like morphologies consistent with images of similar 2D MOFs. Nitrogen
gas sorption isotherms show that FeTHT has a Brunauer-Emmett Teller (BET) surface area
of 441 m
2
/g (Figure 4.3), which is similar to the surface areas reported for analogous cobalt
and iron THT frameworks of 370 and 526 m
2
/g, respectively.
51,55

Figure 4.1. Synthesis and structure of the FeTHT framework. The plot shows a comparison
of the experimental (red; synchrotron radiation with λ = 0.414576 Å) and simulated (black)
PXRD patterns, and the inset shows an SEM image at 75 low magnification.

98

Figure 4.2. Spacefilling model of the fragment used to generate the unit cell of FeTHT.

Figure 4.3. SEM images of FeTHT.

Figure 4.4. Nitrogen isotherms performed on the as-prepared FeTHT at 77 K reveal a BET surface
area of 441 m
2
/g.

Following oxidation of these frameworks in air (3 days at 60 °C), the PXRD pattern of
FeTHT samples show similar prominent peaks at 1.21°, 2.18°, 2.42°, 3.21° compared with
the as-prepared sample (Figures 4.4-4.6), suggesting analogous structural environments in
99

the ab plane after O2-exposure. The broad peak at 7.09° shifts to lower angles and sharpens
to reveal features consistent with the simulated spectrum upon cooling, indicating an
increase in interlayer ordering and a decrease in interlayer distance (from 3.34 Å to 3.23
Å). No significant change in film thickness or morphology was observed by AFM after
oxidation of the framework (Figure 4.7). The measured BET surface area of the oxidized
FeTHT sample was determined to be 27.6 m
2
/g (Figure 4.8), indicating a significant loss
of accessible surface area. A similar reduction in the available surface area from 370 to 50
m
2
/g was previously reported for Fe(tri)2 MOF upon oxidation, where tri = 1,2,3-triazolate,
and was attributed to occupation of the pores by charge-balancing species (BF4
–
).
26

Figure 4.5. Overlay of the variable temperature PXRD patterns of an oxidized FeTHT sample
exposed to air for 3 days at 60 °C. Data was collected at 100 (blue), 200 (green), 295 (orange), and
340 (red) K.

Figure 4.6. Overlay of the variable temperature PXRD patterns of an oxidized FeTHT sample
exposed to air for 3 days at 60 °C, focusing on the [100] (left) and [001] (right) reflections. Data
was collected at 100 (blue), 200 (green), 295 (orange), and 340 (red) K.

100

Figure 4.7. PXRD pattern of an oxidized FeTHT sample exposed to air for 3 days at 60 °C. Data
was collected at 295 K.

Figure 4.8. Atomic Force Microscopy (AFM) studies of an oxidized FeTHT sample exposed to
air for 3 days at 60 °C. The measured film thickness is 275(28) nm.

101

Figure 4.9. Nitrogen isotherms performed on an oxidized FeTHT sample exposed to air for 3 days
at 60 °C. Data collected at 77 K after air-exposure reveal a BET surface area of 27.6 m
2
/g (red:
adsorption, blue: desorption).

4.2.2 Density Functional Theory (DFT) Calculations

To investigate the electronic structure of the FeTHT framework in the absence of ammonium
guest species, first-principles calculations were carried out on the modelled structure using plane-
wave density-functional theory (DFT; see computational methods section for details). The
calculations predict a ground-state magnetization density with one unpaired electron per Fe ion,
consistent with Fe
3+
in a distorted-octahedral crystal field. A strong preference for ferromagnetic
coupling between layers and a weaker frustrated antiferromagnetic coupling within the 2D sheets
produces a net magnetic moment of 2.2 μB per unit cell. The predicted equilibrium geometry is in
good agreement with the experimentally-measured lattice parameters, and similar results were
obtained from DFT+U calculations with a Hubbard U correction of 5 eV applied to the Fe d states.

As in our previous calculations on the analogous CoTHT framework,
51
the electronic band
dispersion suggests bulk FeTHT to be a semi-metal with a small density of metallic states at the
Fermi energy (Figure 4.10). These states correspond to chains of Fe d and S p states, with a strong
interlayer interaction along the c axis, corresponding to the Γ-A segment of the band structure,
producing a large dispersion of ~1.8 eV. Carrier masses ranging from 0.55-6.72 𝑚 𝑒 were
estimated, which are comparable to the values of 0.29-8.04 𝑚 𝑒 predicted for the CoTHT
analogue.
51
The calculated dispersion also contains a small number of metallic states along the M-
Γ segment of the dispersion, corresponding to in-plane conductivity, with a carrier mass around
102

0.99 𝑚 𝑒 . Semi-metallic electronic structures were also predicted for both the FeTHT and CoTHT
frameworks with DFT+U calculations, which provide confidence that these results are not an
artefact of unphysical delocalisation of the Fe d electrons. Prior electronic-structure calculations
on the FeTHT framework with explicit ammonium guest molecules predict a narrow bandgap
(~350 meV),
55
highlighting the sensitivity of the electronic structure to subtle environmental
factors. The predicted elimination of this gap upon removal of the NH 4
+
provides a strong
indication that the modified FeTHT reported here should exhibit distinct electronic behaviour
relative to the previously-reported semiconducting FeTHT.

Figure 4.10. Calculated band dispersion and electronic density of states curves for the FeTHT
framework with no guest species. The blue and red lines denote electronic states in the “up” and
“down” spin channels. The thick black lines mark the regions of the dispersion used to evaluate
𝜕 2
𝐸 (𝑘 ) 𝜕 𝑘 2
⁄ to estimate the carrier effective masses discussed in the text.

Computational studies of the Ni3(HITP)2 framework
43
suggested that a staggered arrangement of
the layers reduced the total energy compared to an eclipsed geometry, whereas our previous
calculations on the CoTHT framework found the latter to be lowest in energy.
51
The calculated
potential-energy surface for stacking faults in FeTHT (Figure 4.11) shows that layer offsets of
~1.0-1.25 Å along the a/b axis or a combined offset of 1.25 Å along both axes are lower in energy
than the eclipsed configuration by ∆𝐸 = 83 and 98 meV per Fe ion, respectively, suggesting that
formation of stacking faults is likely to be somewhat facile in this system, consistent with the broad
[001] PXRD peak observed. This result is also consistent with prior DFT calculations for the
FeTHT framework with ammonium guest species.
55
As for the CoTHT framework, the
103

calculations predict that the stacking faults would reduce the density of states around the Fermi
energy and introduce a gap into the conduction band, both of which are intuitive given the nature
of the metallic states. These stacking faults, however, are also predicted to have a minimal impact
on the lattice parameters and magnetic configuration. These results thus suggest that intrinsic
stacking faults, as well as environmental factors, may play key roles in defining the electronic
structure of FeTHT and inducing a semiconductor-to-metal transition.

Figure 4.11. Calculated potential-energy surface associated with layer offsets (stacking faults) in
the pristine FeTHT framework.

4.2.3 X-ray Photoelectron Spectroscopy (XPS) and Magnetism Studies

The surface composition of the FeTHT films was investigated using X-ray photoelectron
spectroscopy (XPS). Survey scans of pristine FeTHT films reveal the presence of Fe, S, C,
and O. Fitting of the Fe 2p, shown in Figure 4.12a, was primarily based on the Fe 2p
multiplet splitting of Fe3O4 which was used as a reference.
68
The peaks at binding energies
of 708.2, 709.2, and 710.3 eV are assigned to Fe
2+
, and the peaks at 710.1, 711.2, 712.3,
and 713.5 eV are assigned to Fe
3+
(See Table 4.1). The Fe 2p region shows mixed Fe
3+/2+

valency with approximately 52% Fe
2+
and 48% Fe
3+
; indicating the presence of mixed
oxidation states in pristine FeTHT.
23
Figure 4.12b shows peak fitting of the S 2p regions
for the pristine sample. The binding energy difference of the S 2p 3/2 - S 2p1/2 doublet was
fixed to 1.2 eV with intensity ratio 2:1.
69,70

104

Figure 4.12. Fitting of the XPS spectra of pristine FeTHT and samples exposed to ambient
atmosphere for 7 days. (a) Fe 2p region showing mixed valency (Fe
2+/3+
) in pristine FeTHT; after
exposure to air, only the Fe
3+
peak is observed. (b) S 2p region in the pristine FeTHT; after
exposure to air, two new doublets at ~166 eV (pink) and ~168 eV (burgundy) appear, indicating
the formation of dithiolene moieties with different degrees of oxidation.

Table 4.1. Peak parameters to fit Fe
2+
and Fe
3+
multiplets in FeTHT. The reference parameters
are of Fe3O4.
68
Fe3O4 has a mixed Fe
3+/2+
valency with a 2:1 ratio, which is comparable to the
FeTHT and thus, the primary reason for reference in fitting the Fe 2p region. Details of the fitting
can be found in the table below. Peak 1 represents the lowest binding energy peak for both the Fe
2+

and Fe
3+
multiplets. Note that the single Fe
3+
satellite (~716 eV) is not in the table because peak
constraints were not held fixed between samples.

Examination of the same XPS regions in a sample after prolonged air exposure reveals
distinct changes in the Fe 2p and S 2p regions. Following 7 days of air oxidation, the first
peak of the Fe
2+
multiplet disappears and the experimental data can be fit by the Fe
3+

multiplet alone (Figure 4.12a), indicating complete surface oxidation of the metal sites. In
the S 2p region, two new sets of peaks appear at higher binding energies of ~166 eV and
~168 eV, indicating concomitant ligand oxidation (Figure 4.12b). These environments are
attributed to sulfenate and sulfinate moieties, as has previously been reported in a palladium
dithiolene molecular complex.
71
To further substantiate these assignments, high resolution
XPS of the corresponding palladium and iron molecular species have been collected before
and after O2 exposure (Figure 4.13-4.16). The S 2p region of the previously reported
Peak 1 (eV) Peak 2 (eV) ∆E
(peak2-peak1)
Peak 3 (eV) ∆E
(peak3-peak2)
Peak 4 (eV) ∆E
(peak4-peak3)
[FWHM] [FWHM] (eV) [FWHM] (eV) [FWHM] (eV)
Fe 2+
Reference 708.3 [1.2] 41.6 709.3 [1.2] 43.2 1.0 710.4 [1.4] 15.2 1.1
Pristine 708.2 [1.0] 41.8 709.2[1.6] 43.1 1 710.3 [1.6] 15.1 1.1
Fe 3+
Reference 710.2 [1.4] 34.9 711.1 [31.8] 31.8 1.1 712.4 [1.4] 22.6 1.1 713.6 [1.4] 10.8 1.2
Pristine 710.1 [1.6] 35.1 711.2 [1.5] 31.9 1.1 712.3 [1.6] 22.5 1.1 713.5 [1.6] 10.5 1.2
7 Days 710.1 [1.6] 35.1 711.2 [1.5] 31.9 1.1 712.3 [1.6] 22.5 1.1 713.5 [1.6] 10.5 1.2
% % % %
105

palladium coordination polymers confirms that aerobic oxidation results in a new oxidized
sulfur moiety, which appears at ~167 eV. This new feature is assigned as the added presence
of sulfenate/sulfinate moieties, which is confirmed by the crystal structure of the resulting
polymer. Likewise, aerobic oxidation of the analogous iron dithiolene complex reproduces
the additional S 2p features observed for the FeTHT framework, which supports the
assignment of these new features as sulfonate/sulfonate moieties. By examining the area of
the peaks in the O2-exposed FeTHT sample, approximately 49% of the thiolate moieties
appear to be oxidized, with a sulfenate:sulfinate ratio of 1:4. These numbers are similar to
the palladium dithiolene complex, in which half of all thiolates are oxidized, with a
sulfenate/sulfinate ratio of 1:2.3.

Figure 4.13. XPS spectra showing the iron 2p (a) and sulfur 2p (b) features from the
[Febdt2](HNEt3)2 molecular complex

Figure 4.14. XPS spectra showing the iron 2p (a) and sulfur 2p (b) features from the
[Febdt2](HNEt3)2 molecular complex after 3 days of oxidation in air
106

Figure 4.15. XPS spectra showing the palladium 3d (a) and sulfur 2p (b) features of a [Pdbdt2]K2
molecular complex after 30 minutes of air exposure

Figure 4.16. XPS spectra showing the palladium 3d (a) and sulfur 2p (b) features from the
oxidized [Pdbdt2]K2 species, {[K4(thf)4(H2O)2.28][Pd-(O2SC6H4S)1.36(OSC6H4S)0.64]2}n,
which forms a crystalline coordination polymer as reported in Zamora
71

These XPS results thus confirm that both the metal center and the ligand scaffold in FeTHT
are oxidized upon prolonged air exposure. This result is consistent with the predominant
contributions from interacting Fe and S orbitals to the partially-occupied bands predicted
by DFT. DFT also predicts a preference for an offset rather than eclipsed layer stacking,
indicating that interlayer Fe-S interactions may play a role in directing stacking manifolds.
As these are not expected to be strong interactions, based on the predicted ease of
generating stacking faults, the formation of new S-O bonds upon oxidation may facilitate
stronger Fe-O interlayer interactions, consistent with the sharpening of the peak at 7.09° by
synchrotron PXRD. As no such oxidized thiolate moieties were evident in the CoTHT
107

material, we attribute the thickness-dependence of the cobalt system to a slower oxidation
process relative to the analogous iron-containing framework.
51

To quantify the bulk composition of the iron oxidation states in FeTHT, magnetic
measurements were performed. It has been previously reported that mixed-valency in
MOFs can enhance conductivity by promoting intervalent charge transfer.
72
Recent studies
of iron-based 3D MOFs have shown that controlling mixed-valency at iron sites can have
dramatic effects on the bulk transport properties, increasing conductivity by as many as
eight orders of magnitude.
25,26,73
By examining the saturation of unpaired spins in the
hysteresis loop (Figure 4.17a), it is estimated that 33% of the iron centers are Fe
3+
, with the
remaining 67% Fe
2+
. These results are in agreement with the XPS studies, as similar mixed
Fe
3+/2+
valency is observed in the as-prepared FeTHT material, albeit a higher degree of
oxidation is estimated by XPS (1:1 from XPS studies vs 2:1 from magnetic studies). The
higher degree of oxidation observed by XPS suggests that the oxidized species are
concentrated towards the surface of the sample, as magnetic and XPS measurements
analyse the bulk vs. surface composition, respectively. After O2-treatment of the FeTHT
sample for 3 days at 60 °C, the observed magnetic moment of the sample increases, which
is consistent with oxidation to give an increased amount of Fe
3+
(Figure 4.17b). This
increased moment reflects an increase of the Fe
2+
/Fe
3+
ratio from 2:1 to 1:1. The magnetic
studies thus indicate a slightly lower degree of oxidation of FeTHT compared to the results
from the XPS studies (only Fe
3+
is detected by XPS upon exposure of the sample to ambient
atmosphere for 7 days). These results demonstrate that O2-treatment is a promising tool for
surface-localized electronic structure modification of 2D MOFs. As numerous device
applications require surface-selective treatment, this accessible methodology is amenable
to the preparation of one-component junctions with chemically-controlled transport
properties. Such devices would enable otherwise challenging sensitivity to temperature
and/or guest species inclusion, as a semiconducting-to-metallic transition on the device
surface could produce switching behaviour from a metal-metal to a metal-semiconductor
junction.
108

Figure 4.17. Magnetic hysteresis loops in pristine FeTHT (a) and samples exposed to air for three
days (b).

Figure 4.18. Magnetic studies of pristine FeTHT. (a) temperature-dependent susceptibility, (b)
Curie-Weiss fit of the high-temperature region (200–300 K), and (c) hysteresis loop collected at 2
K.

4.2.4 Resistivity Studies

The temperature-dependent resistivity of FeTHT was measured using the four-point Van
der Pauw geometry. Films of thicknesses ranging from 84(8) nm to 410(41) nm, as
determined by AFM studies (Figure 4.19), were analysed under identical conditions. Ohmic
contacts, as demonstrated by linear I-V traces (Figure 4.20), were made using graphitic
carbon adhesive in ambient atmosphere with minimal (<10 min) air exposure. At 300 K,
the measured bulk resistivity of the 275 nm-thick FeTHT film is 5.4 Ω·cm, which
corresponds to a conductivity of 0.2 S/cm (Figures 4.21 and 4.27, and Table 4.2). As the
sample is cooled from 300 K to 140 K, the film displays an increase in resistivity, as
109

expected for a semiconducting material where transport is dominated by thermally-
populated carriers (Figures 4.21 and 4.27). This behaviour is similar to that observed for
films and pressed pellets of CoTHT at elevated temperatures, and is consistent with the
prior report of semiconducting behaviour for Fe3(THT)2(NH4)3.
51,55
An Arrhenius fit to the
resistivity data recorded between 300 and 200 K indicates an activation energy of 12.8 meV
for carrier promotion (Figures 4.22 and 4.23). This small activation barrier suggests that
defects, most likely grain boundaries in the 2D sheets, play a significant role in the transport
properties of this material.

Figure 4.19. Typical Atomic Force Microscopy (AFM) studies of FeTHT films. The measured
thickness of the FeTHT film is 109(11) nm.

Figure 4.20. Variable-temperature I-V traces of FeTHT demonstrating Ohmic behavior.

110

Figure 4.21. Overlay of the temperature-dependent resistivity data for FeTHT films with
thicknesses ranging from 84(8) to 176(18) nm.

Figure 4.22. Arrhenius plots for FeTHT films with thicknesses of 84(8) nm (top left), 97(10) nm
(top right), 109(11) nm (bottom left), and 176(18) nm (bottom right). The data was fit from 300 –
230 K.

111

Recent DFT studies and a related examination of the temperature-dependent resistivity data
for the Ni3(HITP)2 framework have suggested that the Mott variable-range hopping model
best describes the conductivity behaviour of this material.
60
Fitting the high-temperature
data for FeTHT to the Mott variable-range hopping model and generating an Arrhenius
plot gives an activation energy for charge hopping ranging between 6.2 and 8.9 meV
(Figure 4.23; Table 4.2). Interestingly, these values are similar to the 6 meV hopping barrier
reported for the Ni3(HITP)2 2D framework. This small activation barrier suggests that
defects, likely grain boundaries in the 2D sheets, play a significant role in the transport
properties of these 2D MOFs.

Figure 4.23. Arrhenius plots using the variable range hopping model for FeTHT films with
thicknesses of 84(8) nm (top left), 97(10) nm (top right), 109(11) nm (bottom left), and 176(18)
nm (bottom right). The data was fit from 300 – 230 K.

As the temperature decreases below 140 K, the FeTHT film undergoes a complex transition
from semiconducting-to-metallic behaviour, resulting in a decreasing resistivity upon
cooling below 100 K (Figures 4.21 and 4.27, and Table 4.2). This behaviour at low
temperatures is as expected for metallic conductivity, where scattering of the carriers is
dominated by lattice vibrations. These observations are similar to the behaviour observed
112

for CoTHT,
51
and are fully-reversible with no hysteresis. Given the complex nature of the
transition, Tmetallic is defined as the lowest transition temperature at which FeTHT becomes
metallic (100 K). No morphological changes are observed by scanning electron microscopy
(Figure 4.24). Variable-temperature synchrotron PXRD studies performed between 300
and 100 K indicate a subtle structural distortion related to a minor contraction of the
interlayer stacking (Figures 4.25 and 4.26). An analogous minor contraction (0.02 Å) was
observed for the CoTHT system.
51
The measured semiconductor-to-metal transition
temperatures (Tmetallic) for FeTHT do not display any appreciable shifts with film thickness
(Figure 4.21 and Table 4.2), in contrast to the CoTHT system where higher semiconductor-
to-metal transition temperatures were observed in thinner films.
51
Our new understanding
of the role of chemical oxidation suggests that the previously-reported thickness
dependence on the conductivity of CoTHT represents an artifact of surface-localized
chemical oxidation. This further suggests that similar control over the transition
temperature in the cobalt material may be brought about through post-synthetic oxidation.

Figure 4.24. SEM images of FeTHT after resistivity studies. (left) medium magnification (1000)
and (right) high magnification (60,000).

Entry Thickness
(nm)
Conductivity (S/cm) @
300 K
Ea
[a]

(meV)
Ea
[b]

(meV)
TMetallic
(K)
1 84(8) 0.3 12.8 8.9 160
2 97(10) 0.1 11.3 7.5 190
3 109(11) 0.1 12.7 7.8 160
4 176(18) 0.03 9.4 6.2 160
5 410(41) 0.02 – 8.5 170
Table 4.2. Summary of electrical transport data for FeTHT films.

[a]
fitted using the Arrhenius equation;
[b]
fitted using the variable range hopping model.

113

Figure 4.25. Overlay of the variable temperature PXRD patterns of pristine FeTHT collected at
100 (blue), 200 (green), 280 (orange), and 295 (red) K.

Figure 4.26. Overlay of the variable temperature PXRD patterns of pristine FeTHT focusing on
the [100] (left) and [001] (right) reflections. Data was collected at 100 (blue), 200 (green), 280
(orange), and 295 (red) K.

Air Exposure Conductivity (S/cm)
T
metallic
(K)
as-prepared 0.2 100
2.5 days 0.02 135
1 month
1.4x10
-3

265
1.5 months
8.0x10
-4

300
Table 4.3. Transport data for an FeTHT sample with 275(28) nm thickness after exposure to
ambient atmosphere for varying times (from Figure 2)

114

To demonstrate the value of post-synthetic modification in determining transport properties
of these 2D materials, resistivity studies were conducted on oxidized samples of FeTHT.
The previously-analysed 275(28) nm-thick FeTHT film was oxidized under ambient
atmosphere, without breaking the contacts, and temperature-dependent resistivity data were
recollected. Following 2.5 days of air exposure, an approximately 10-fold increase in the
resistivity is observed, from 5.2 to 47.5 Ω·cm at 300 K (Figure 4.27). Remarkably, this
coincides with an increase in Tmetallic from 100 to 135 K. Further oxidation of the FeTHT
sample for a total of 1 month led to an increase in Tmetallic to 300 K with a corresponding
increase in resistivity to 735 Ω·cm. After a total exposure of 1.5 months, the sample
displays pure metallic conductivity at 300 K (ρ = 1,260 Ω·cm) while maintaining a
transition to semiconducting behavior at higher temperatures. The development of room-
temperature metallic conductivity can be accelerated by heating a 278(28) nm-thick
FeTHT sample in air at 60 °C for 3 days (Figure 4.28). To isolate the influence of O2 on
this shift in the semiconducting-to-metallic transition temperature, we have performed
control studies under an inert atmosphere. Heating the sample at 60 °C for 3 days under 1
atm of helium shows a negligible change in Tmetallic (250 K) (Figure 4.29), confirming that
chemical oxidation with O2 is responsible for the observed shift in Tmetallic. This result
represents the first demonstration of a tunable semiconducting-to-metallic transition in a
MOF through post-synthetic chemical modification.

Figure 4.27. Temperature-dependent resistivity data for FeTHT films with 275(28) nm thickness,
as-prepared with minimal (<10 min) air exposure (blue) and following exposure to ambient
atmosphere for 2.5 days (red), 1 month (green) and 1.5 months (orange).
115

Figure 4.28. Temperature-dependent resistivity data for a FeTHT film with a thickness of
275(28) nm as-prepared (blue) and after 3 days of air exposure at 60°C (red).

Figure 4.29. Temperature-dependent resistivity data for a FeTHT film with a thickness of
275(28) nm after 3 days at 60°C under 1 atmosphere of helium.

4.3 Conclusions

We report here a modified synthesis for the recently-reported iron 2,3,6,7,10,11-
tripheylenehexathiolate (FeTHT) MOF, generating a material with unprecedented electronic
behaviour. The peaks observed in the synchrotron PXRD pattern are similar to those of the CoTHT
framework, suggesting analogous structural environments for the cobalt and iron systems.
Temperature-dependent resistivity studies reveal a transition from semiconducting to metallic
upon cooling, which is analogous to the behaviour observed for CoTHT. This semiconducting-to-
metallic character of the FeTHT system investigated here is in contrast with the previously
116

reported Fe3(THT)2(NH4)3 framework, which exhibits only semiconducting behaviour, and is
attributed to the elimination of NH4
+
guest species through the modified synthesis as predicted by
DFT calculations. Interestingly, ambient chemical oxidation of this material leads to a positive
shift in the semiconducting-to-metallic transition temperature with time, resulting in a material
with metallic character at room temperature. Through a combination of X-ray photoelectron
spectroscopy (XPS) and magnetic studies, we identified that both the metal center and the ligand
scaffold in FeTHT are oxidized upon O2-treatment. These results demonstrate, for the first time,
a room-temperature semiconductor-to-metallic transition in a MOF induced by post-synthetic
chemical. The surface-localized electronic structure modification by O2-treatment is expected to
enable single-component junctions with thermal and chemical sensitivity with potential
applications in chemiresistive sensing and optoelectronics.

4.4 Experimental Details

4.4.1 General Considerations

All manipulations of air and moisture sensitive materials were conducted under a nitrogen
atmosphere in a Vacuum Atmospheres drybox or on a dual manifold Schlenk line. The glassware
was oven-dried prior to use. Acetonitrile and dichloromethane were degassed with nitrogen and
passed through activated alumina columns and stored over 4 Å Linde-type molecular sieves. Ethyl
acetate, water, and ethanol were placed under vacuum and refilled with nitrogen (10 ). Deuterated
solvents were dried over 4 Å Linde-type molecular sieves prior to use. Elemental analyses were
performed by Robertson Microlit Laboratories, 1705 U.S. Highway 46, Suite 1D, Ledgewood,
New Jersey, 07852. All the chemical regents were purchased from commercial vendors and used
without further purification. The ligand 2,3,6,7,10,11-triphenylene-hexathiol
74
(THT) was
prepared according to the reported procedures. Water was deionized with the Millipore Milli-Q
Synergy system (18.2 M·cm resistivity). All other chemical regents were purchased from
commercial vendors and used without further purification.

4.4.2 Synthesis of FeTHT

The iron triphenylene-2,3,6,7,10,11-hexathiolate framework, FeTHT, was prepared according to
the reported procedure for the analogous cobalt framework.
51,67
A 120 mL glass jar was charged
117

with an aqueous solution of FeCl2·4H2O (40 mg, 0.21 mmol, 5 mM, 40 mL volume). Separately,
a suspension of triphenylene-2,3,6,7,10,11-hexathiol (THT) (2 mg, 0.005 mmol) in N-Methyl-2-
pyrrolidone (NMP) (0.2 mL) was diluted with ethyl acetate until the total volume of the suspension
reached 20 mL, sealed, and briefly sonicated to form an uniform suspension. Ethyl acetate (40 mL)
was gently added to the aqueous solution of iron(II) chloride to create a liquid-liquid interface.
The suspension of THT and NMP was then gently added to the ethyl acetate layer and the jar was
sealed and allowed to stand. A black film appeared at the liquid-liquid interface over 5 days. The
film was deposited onto glass supports by pulling the substrate through the film. The deposited
films were subsequently washed with water and allowed to evaporate to dryness. Alternatively,
the black solid of FeTHT was collected by filtration and washed with water and methanol for bulk
powder analyses.

FeTHT powders or films were oxidized by placing them in an oven at 60 °C for 3 days. Anal.
Calcd for Fe3(THTox)2 • 3MeOH • 12H2O (Fe3C45H48S12O44), where THTox = C18H6O9S6: C,
32.38; H, 2.62; Fe, 9.08; S, 23.05. Found: C, 32.89; H, 2.86; Fe, 9.28; S, 23.53.

Powder X-ray diffraction (PXRD) studies were performed on a Rigaku Ultima IV X-Ray
diffractometer in reflectance parallel beam/parallel slit alignment geometry. The measurement
employed Cu K line focused radiation at 1760 W (40 kV, 44 mA) power and a Ge crystal detector
fitted with a 2 mm radiation entrance slit. Samples were mounted on zero-background sample
holders and were observed using a 0.08° 2 step scan from 2.0 – 40.0° with an exposure time of
0.4 s per step. No peaks could be resolved from the baseline for 2 > 35°.

High resolution synchrotron powder X-ray diffraction data was collected using the 11-BM
beamline mail-in program at the Advanced Photon Source (APS), Argonne National Laboratory,
with an average wavelength of 0.414576 Å. Discrete detectors covering an angular range from 0.5
to 30º 2θ are scanned over a 34º 2θ range, with data points collected every 0.001º 2θ and scan
speed of 0.01º/s. An Oxford Cryosystems Cryostream Plus device allowed for sample temperatures
to be controlled over a range of 100-340 K.

118

Gas Sorption Measurements

Brunauer-Emmett-Teller (BET) measurements were performed on a Nova 2200e surface area and
pore size analyzer (Quantachrome Instruments, Inc.). Samples were degassed for 3 days at 60 °C
in vacuo prior to measurements.

Modeling

Molecular modeling of FeTHT was carried out using the Materials Studio (version 8.0) suite of
programs by Accelrys. The molecular fragment used to generate the model is shown in Figure 4.2.
The unit cell was constructed starting with a primitive hexagonal unit cell with space group
P6/mmm using cell parameters a = b = 22.52 Å and c = 3.3 Å. The structure was optimized with
Materials Studio Forcite calculations using geometry optimization and universal forcefield
methods. The MS Reflex module was used to calculate the expected PXRD patterns. Line
broadening for crystallite size was not calculated. Comparison of the simulated and experimental
PXRD patterns verified the simulated structure.

Scanning Electron Microscopy (SEM)

SEM images of FeTHT on glass substrates were collected using a JEOL-7001F or FEI Nova
NanoSEM 450 operating at 10 or 15 kV with 5 nA of probe current.

Conductivity Measurements

FeTHT films were deposited onto glass supports by pulling the substrate (glass) through the
FeTHT film formed at the liquid-liquid interface, as described previously. The deposited films
were subsequently washed with water and allowed to evaporate to dryness. Conductivity
measurements were performed using a custom set up integrated into a 14T Quantum Design
Dynacool Physical Properties Measurement System. A Keithley 6220 Precision Current Source
(excitation currents of 1-50 nA) was used to trigger and control a Keithley 2182A nanovoltmeter.
In order to minimize errors associated with contact resistance and drift voltages, a Keithley 2172
matrix switch equipped with a Keithley 6536 Hall effect card was used to alternate the direction
of the applied current. Because of difficulty associated with preparing samples with uniform
dimensions, all measurements were performed in a four-point probe Van der Pauw geometry.
119

Copper wire contacts were attached to the films using conductive carbon paint and soldered onto
a Quantum Design puck with resistivity option. All measurements were performed under a reduced
pressure of ~10 torr.

Magnetic Studies

Temperature-dependent susceptibility measurements of pristine FeTHT show a paramagnetic
response down to 2 K, as illustrated in Figure 4.18a. A fit of the high temperature magnetic
susceptibility (175–275 K) to the Curie-Weiss law χ = C/(T-Ɵcw) yields an effective paramagnetic
moment, μeff = 2.20 μB

per formula unit and a Curie-Weiss temperature Ɵcw of 0.60 K. The fit,
shown in Figure 4.18b, shows a slight positive deviation corresponding to the positive Ɵcw of 0.60
K, which indicates that the dominant exchange interaction among the Fe ions is slightly
ferromagnetic. During magnetization, the sample does not saturate up to 14 T (Figure 4.18c),
which is characteristic of a paramagnetic material.

Atomic Force Microscopy (AFM)

FeTHT films were deposited onto glass supports by pulling the substrate (glass) through the
FeTHT film formed at the liquid-liquid interface, as described previously. The deposited films
were subsequently washed with water and allowed to evaporate to dryness. AFM topography
images were collected in tapping mode using an Agilent 5420 SPM instrument operating in tapping
mode. The probe tips were Tap300-G Silicon AFM probes (resonant frequency 300 kHz, force
constant 40 N/m) purchased from Budgetsensors.com and aligned prior to use. Images were
collected with a scan rate of 0.1 lines per second and over an area of 40 µm. All samples were
imaged under one atmosphere of air at room temperature.

X-Ray Photoelectron Spectroscopy (XPS)

FeTHT samples were mounted and transferred via sealed vacuum suitcase to a lab-based
monochromatic Al-Kα with a hemispherical analyzer located at the Analytical and Diagnostics
Laboratory at Binghamton University. Measurements were performed with a pass energy of 23.5
eV, corresponding to an instrumental resolution of 0.51 eV, determined from analyzing the Fermi
edge and Au 4f7/2 peak of gold foil references. Lorentzian broadening of Fe 2p multiplet peaks,
120

necessary to fit Voight profiles in peak fitting, was approximated from the instrumental resolution
and FWHM of Fe3O4 peaks.
68
Energy calibration of all core regions were made to Au 4f7/2
peak(84.0 eV) of gold foil references.

4.4.3 Computational Modelling

Periodic density-functional theory (DFT) calculations on the FeTHT framework were carried out
within the pseudopotential plane-wave formalism implemented in the Vienna Ab initio Simulation
Package (VASP) code.
75

Electron exchange and correlation were modelled with the PBEsol functional
76
with the DFT-D3
dispersion correction
77
(i.e. PBEsol-D3). A subset of calculations were repeated with a Hubbard
correction of 𝑈 eff
= 5 eV applied to the Fe d states using the Dudarev method.
78

The ion cores were modelled using projector augmented-wave (PAW) pseudopotentials
79,80
with
the H 1s, C 2s/2p, S 3s/3p and Fe 4s, 3d and 3p electrons included in the valence shells. Based on
our previous studies of the analogous Co framework,
51
an 800 eV kinetic-energy cutoff was
employed for the plane-wave basis and the electronic wavefunctions were modelled using a Γ-
centred Monkhorst-Pack k-point mesh
81
with 1×1×5 subdivisions, reduced appropriately for
supercell calculations. A Gaussian smearing of 0.01 eV was used to determine partial band
occupations.

Tolerances of 10
-8
eV and 10
-2
eV Å
-1
were applied to the total energy and forces during
minimisation of the electronic wavefunctions and geometry optimisations, respectively. The
precision of the charge-density grids was set to avoid aliasing errors and the PAW projection was
performed in real space.

The electronic density-of-states (DoS) curves were evaluated from a single-point calculation with
an increased k-point density of 2×2×15 subdivisions and a larger smearing width of 0.05 eV. The
charge density from these calculations was then used to model band dispersions by calculating the
eigenvalues on strings of k-points along high-symmetry directions in the Brillouin zone non self-
consistently.
121

Equilibrium Geometry and Magnetic Structure

To identify the most energetically-favourable magnetic configuration, we made a starting
assumption of a 50/50 mixture of Fe
2+
and Fe
3+
ions in a square-planar crystal field with zero and
one unpaired electron, respectively. A 50/50 mixture of charge states is not compatible with the
three Fe ions in the crystallographic primitive unit cell of the FeTHT framework, so we
constructed trial magnetic unit cells based on 1×1×2 and 2×2×1 expansions of the primitive cell
with 6/12 Fe ions and enumerated all symmetry-inequivalent arrangements of the Fe
2+
and Fe
3+

ions using the Transformer code,
82
yielding in a total of three and 30 initial configurations in the
two cells, respectively. A single-point energy calculation was then carried out on each model, with
the magnetic moments allowed to relax during the optimisation of the electronic wavefunctions.
We note that this is a more systematic approach than that taken in our previous study on the cobalt
analogue (CoTHT).
51

All 33 initial models relaxed to configurations with magnetic moments around ±2 BM per Fe ion,
consistent with all the ions adopting the Fe
3+
oxidation state. The calculations on the 1×1×2
supercell produced one ferromagnetic and one frustrated antiferromagnetic configuration with
total moments of 6.3 and 2.0 BM per formula unit (i.e. per single unit cell), respectively. The latter
configuration is calculated to be 20 meV per Fe ion lower in energy than the fully ferromagnetic
arrangement and is based on chains of Fe ions with parallel spin along the c axis and a mix of
ferromagnetic and antiferromagnetic coupling between chains within the layers. A configuration
constrained to have ferromagnetic interlayer and antiferromagnetic intralayer coupling was found
to be ~0.1 eV per ion higher in energy than the fully ferromagnetic configuration, indicating a
strong preference for ferromagnetic coupling between layers.

The majority of the 30 2×2×1 models relaxed to antiferromagnetic ground states with an equal
number of both spin states, but a small number adopted configurations with two and four excess
spins and net magnetic moments of 1.1 and 2.2 BM per formula unit, respectively. The energetic
differences were on the order of 1-2 meV per Fe ion, suggesting weak intralayer coupling.

122

Three of the initial screened configurations were selected for full geometry optimisation, viz. the
frustrated antiferromagnetic 1×1×2 expansion (Model 1), the antiferromagnetic 2×2×1 expansion
(Model 2), and the 2×2×1 expansion with the larger magnetic moment of 2.2 BM per formula unit
(Model 3). The optimised lattice parameters, total magnetic moments and total energies are
compared in Table S1.

Expt. Model 1 Model 2 Model 3
Supercell - 1×1×2 2×2×1 2×2×1
a [Å] 22.52 23.23 23.22 23.22
c [Å] 3.3 3.299 3.256 3.335
c/a 0.147 0.142 0.143 0.144
V [Å
3
] 1449 1541 1553 1557
M [BM] 1.87 3.04 -0.03 2.21
E0 [eV per Fe ion] - -155.24 -155.30 -155.30
Table 4.4. Optimised lattice parameters, total magnetic moments and total energies of three trial
magnetic supercells of the FeTHT framework. The lattice parameters and magnetic moments are
given with respect to the crystallographic unit cell (i.e. containing three Fe ions), and the total
energies are given per Fe ion. The experimental values are shown in the second column for
comparison.

Given the weak intralayer coupling, the 1×1×2 and net magnetic 2×2×1 configurations (Models
1 and 3) should be roughly equivalent, and the differences in Table 1 provide an estimate of how
much variation can be expected due to differences in the choice of unit cell and associated technical
parameters such as the k-point sampling. With this in mind, the only major difference between all
three models is the magnetic moment. Since Models 1 and 3 can be represented to some level of
approximation with a single crystallographic unit cell, which is much more computationally
tractable than either supercell, we opted to do this in our production calculations. This also has the
advantage of making the calculations directly comparable to those in our previous work on the
CoTHT analogue.
51

Using the single-cell configuration, we obtained optimised lattice parameters of a = 23.21 and c =
3.334 Å, which are a good match to the experimental measurements of 22.52 and 3.3 Å
123

respectively. The calculated spin density (Figure 4.30) shows that the unpaired electrons are
localised to the Fe d orbitals, as expected, with a small amount of density on the coordinating S
atoms and very little on the ligand 𝜋 system.

Figure 4.30. Spin density of the lowest-energy magnetic configuration of the FeTHT
framework.

With the more accurate convergence settings used in the electronic-structure calculations, we
obtain a magnetic moment of 2.15 BM per formula unit, which is very close to the low-temperature
experimental measurement of 2.2 BM, albeit with a different origin. We also note that the weak
intralayer coupling predicted by the calculations would, in principle, allow the system to adopt
alternative spin configurations at finite temperature or in response to perturbations to the crystal
field, for example oxidation, which is again consistent with experimental observations.

Electronic Structure

Figure 4.31 shows the band dispersion and electronic density of states (DoS) curves calculated for
the FeTHT framework. As for the CoTHT analogue, the calculations predict this system to be a
semi-metal, with a cluster of bands crossing the Fermi energy along the Γ-A direction in reciprocal
space. This corresponds to the c direction in real space, and an orbital-density plot of the states
within 25 meV of the Fermi energy (Figure 4.32) shows that that the partially-occupied bands in
124

both frameworks correspond primarily to chains of interacting Fe and S orbitals with a small
contribution from the ligand π system.

As in our previous study, bands crossing the Fermi energy were identified and the one-dimensional
𝐸 (𝑘 ) dispersion relation fitted to a quadratic function to estimate the carrier effective masses
according to:
51

1
𝑚 ∗
=
1
ℏ
2
𝜕 2
𝐸 (𝑘 )
𝜕 𝑘 2

(1)

Figure 4.31. Calculated band dispersion and electronic density of states curves for the FeTHT
framework. The blue and red lines denote electronic states in the “up” and “down” spin channels.
The thick black lines indicate the parts of the dispersion used to evaluate 𝜕 2
𝐸 (𝑘 ) 𝜕 𝑘 2
⁄ for
estimating the carrier effective masses.

125

Figure 4.32. Orbital density showing states within 25 meV of the Fermi energy in the FeTHT (a)
and analagous CoTHT (b) framework studied in our previous work (Ref.
51
).

The relatively large bandwidth of up to ~1.8 eV results in low carrier effective masses of -0.55 to
-6.72 and 1.36 to 2.78 𝑚 𝑒 , which can be compared to masses of -0.42 to -1.52 and 0.29 to 8.04 𝑚 𝑒
in CoTHT.
51
Some of the bands along the M- Γ segment, corresponding to in-plane conductivity,
are also metallic, with a carrier mass around 0.99.

Stacking Faults

Previous studies on the layered Ni3(HITP)2 framework (HITP = 2,3,6,7,10,11-
hexaiminotriphenylenesemiquinonate) showed this system to preferentially adopt a staggered
layer arrangement.
43
In contrast, however, our previous study on CoTHT found that an eclipsed
layer configuration was the most energetically favourable.
51

126

To examine the layer stacking in the FeTHT framework, we carried out a similar study to explore
the potential-energy surface associated with the relative positions of the layers in a bilayer 1×1×2
supercell expansion (Figure 4.33). As in our previous work, a series of single-point calculations
were performed with one layer displaced relative to the other by up to 4 Å along the a and b axes,
and at each (∆a, ∆b) displacement coordinate the energies of interlayer spacings from 3.084 to
3.584 Å were calculated and the minimum obtained from a polynomial fit (see Figure 4.34).

Figure 4.33. Calculated potential-energy surface associated with layer offsets in the FeTHT
framework.

In contrast to the Co analogue, the fully eclipsed configuration of the FeTHT framework is 0.21
eV (34 meV per Fe ion) higher in energy than the global minima, which correspond to
displacements along the a or b directions of ∆ ≈ 1.0-1.25 Å. There is also a secondary local
minimum located close to ∆ = (1.25, 1.25) which is a very small 14 meV (2.4 meV per Fe ion)
higher in energy. To investigate further, the two global minima and the low-lying local minimum
were fully volume relaxed. Figure 4.35 shows a view along the c axis of each of the three staggered
bilayer structures, and the lattice parameters, magnetic moments and total energies are compared
to those of an eclipsed bilayer configuration in Table 4.4.

127

Figure 4.34. Illustration of the procedure for locating the minimum-energy c-axis length
(interlayer spacing) at a layer displacement of ∆a = 1.25 and ∆b = 1.25 Å. The blue markers show
the calculated total energies, the red line shows a 1D cubic spline interpolation through the data,
and the cross marks the position of the identified energy minimum.

Expt. Eclipsed (1.00, 0.00) (0.00, 1.00) (1.25, 0.25)
a [Å] 22.52 23.210 23.226 23.226 23.218
c [Å] 3.3 3.334 3.159 3.159 3.163
c/a 0.147 0.150 0.136 0.136 0.136
V [Å
3
] 1449 1556 1476 1475 1476
M [BM] 1.87 2.17
a
2.11 2.11 1.71
E0 [eV per Fe
ion]
- -155.30
a
-155.39 -155.39 -155.40
Table 4.5. Optimised lattice parameters, total magnetic moments and total energies of the FeTHT
framework with eclipsed layers and the three different staggered layer configurations (stacking
faults) shown in Figure 4.35. As in Table 4.4, The lattice parameters and magnetic moments are
given with respect to the crystallographic unit cell, and the total energies are given per Fe ion. The
experimental values are shown in the second column for comparison.
a
Values computed for a
1×1×2 expansion of the optimised single-cell model.

After optimisation, the staggered minima are further lowered in energy to 83 meV per Fe ion
relative to the eclipsed configuration. The structure displaced along both the a and b axes (Figure
4.35c) becomes the global minimum, being a comparatively small 15 meV per Fe ion lower in
energy than the two structures displaced along one of the in-plane directions. All three staggered
configurations result in a slight expansion of the a and b axes by 0.01-0.02 Å and a contraction of
the c axis by 0.17-0.18 Å (5.2 %), which together produce a 5 % reduction in the unit-cell volume.
128

The layer offsets do not appear to change the magnetic ordering, but the calculations predict that
simultaneous displacement along both a and b axes could reduce the total magnetic moment by up
to ~20 %.

Figure 4.35. Optimized structures of bilayer models of FeTHT with layer displacements
corresponding to the local minima in Fig. 4.32, viz. ∆ = (1.00, 0.00) (a), ∆ = (0.00, 1.00) (b) and ∆
= (1.25, 1.25) (c). For clarity, the carbon skeletons of the THT ligand in the two layers are coloured
red and blue.
129

The layer displacement has a small effect on the calculated electronic structure (Figure
4.36), in particular opening a gap in the conduction band at ~1 eV above the Fermi energy, which
is similar to the effect of increasing the interlayer spacing in the CoTHT analogue.
51
To quantify
the changes in the density of states around the Fermi energy, we estimated the room-temperature
(𝑇 = 300 K) concentration of conduction electrons according to:

Figure 4.36. Electronic density of states (DoS) curves of bilayer models of FeTHT in the eclipsed
configuration (a) and staggered configurations corresponding to displacements along the a and b
axes of ∆ = (1.00, 0.00) (b), ∆ = (0.00, 1.00) (c) and ∆ = (1.25, 1.25) (d). The positive and negative
curves on each subplot denote the DoS in the “up” (blue) and “down” (red) spin channels.

𝑛 𝑒 (𝑇 ) = ∫ 𝑓 (𝐸 )𝑔 (𝐸 )𝑑𝐸 𝐸 max
𝐸 F
= ∫
1
1 + exp (
𝐸 − 𝐸 𝐹 𝑘 B
𝑇 )
𝑔 (𝐸 )𝑑𝐸 𝐸 max
𝐸 F

(2)

130

where 𝑓 (𝐸 ) is the Fermi-Dirac distribution, 𝑔 (𝐸 ) is the electronic density of states, 𝑘 B
is the
Boltzmann constant and the integral runs from the Fermi energy to the highest-energy conduction
states in the calculations. We obtained a value of 8.8 × 10
25
m
-3
for the eclipsed structure, which,
compared to a typical value of 10
28
m
-3
for a “good” metal, supports the predicted semi-metallic
nature of the FeTHT system. The calculated concentrations of 7.2 × 10
25
and 7.5 × 10
25
m
-3
for
the structures with layers displaced along the a/b and both axes, respectively, suggest that layer
misalignment reduces the density of conduction states close to the Fermi energy, as would be
expected given the nature of the conductive states (c.f. 4.32). This result is again similar to the
behaviour of the CoTHT analogue.
51

These results indicate that the FeTHT system may show a propensity for layer misalignment, in
contrast to the CoTHT analogue, and the effect of this, together with environmental factors such
as the presence of guest molecules in the pores, may play a role in the measured temperature
dependence of the resistivity.

Effect of a Hubbard U Correction

Systems with strongly-localised electrons, such as those containing d- and f-block elements, can
pose a challenge to (semi-)local DFT methods such as PBEsol, where self-interaction error tends
to overly delocalise the electrons and predict unrealistic electron energies and other calculated
physical properties such as the magnetic ground state.
83
Of particular importance to this work is
that in some systems the error can lead to calculations predicting insulators to be metals.
84
A
straightforward and computationally-efficient method to correct for self-interaction is to apply a
Hubbard U correction to specific atomic states and introduce an energy penalty to force the orbital
occupations towards integer values.
78,83

To test the effect of a Hubbard correction on our results, we performed an identical set of
calculations on the single-cell model of the FeTHT framework with a Hubbard correction of U =
5 eV applied to the Fe d states. This resulted in a slight lengthening of the a and c lattice constants
by ~0.1 Å and a 1 % increase in the cell volume (Table 4.5). The frustrated antiferromagnetic
configuration holds, but stronger localisation of the Fe d electrons increases the total magnetic
moment from 2.2 to 2.4 BM per F.U., with the individual Fe moments increasing from ±2.0 to
131

±2.8 BM per ion. The correction leads to noticeable changes in the electronic structure (Figure
4.37), but the semi-metallic nature of the framework is preserved, and the carrier effective masses
obtained from the curvature of bands along the M-Γ and Γ-A directions are of comparable
magnitude to the bare DFT values, albeit with larger variation and heavier extremes.

Similar DFT+U calculations on the CoTHT analogue with an identical correction of U = 5 eV
applied to the Co d states gave similar results (Table 4.6, Figure 4.38), leading to expansion of the
c axis by ~0.3 Å, a 9.5 % increase in the cell volume, and an increase in the total and Co magnetic
moments from 2.2 to 4.8 BM per F.U. and 2.0 to 2.8 BM per Co ion, respectively. The larger
changes compared to the FeTHT framework could be due to the predicted fully-ferromagnetic
ground state in the Co framework, although systematic studies on transition-metal oxides suggest
that the Co system may not need as large a U value as the Fe one.
85
Once again, despite noticeable
changes to the electronic band dispersion and DoS, the predicted semi-metallic electronic structure
is preserved, and the calculated carrier effective masses are of a similar magnitude.

These calculations provide some confidence that the predicted semi-metallic electronic structure
of the pristine framework is not an artefact from our use of the semi-local PBEsol exchange-
correlation functional, although we note that more sophisticated treatments of electron correlation
may produce a different picture. However, we consider such an undertaking to be beyond the scope
of this investigation.

Expt.
Calc.
PBEsol-D3 PBEsol-D3 + U
a [Å] 22.52 23.210 23.308
c [Å] 3.3 3.334 3.342
c/a 0.147 0.150 0.143
V [Å
3
] 1449 1556 1572
M [BM] 2.20
2.15
(1.95, 1.95, -
1.95)
2.43
(2.83, 2.83, -
2.83)
Table 4.6. Calculated lattice parameters and magnetic moments of the FeTHT framework
structure obtained using the PBEsol-D3 functional with and without a Hubbard correction of U =
5 eV applied to the Fe d states. For each calculation, both the total magnetic moment and the
moments of the three Fe ions are given. Experimental values are shown in the second column for
comparison.
132

Figure 4.37. Band dispersion and electronic density of states (DoS) curves of the FeTHT
framework calculated with PBEsol (top) and PBEsol+U with a Hubbard correction of U = 5 eV
applied to the Fe d states (bottom). As in Figure 4.31, the blue and red colours denote electronic
states in the two spin channels, and the thick black lines indicate parts of the dispersion used to
evaluate 𝜕 2
𝐸 (𝑘 ) 𝜕 𝑘 2
⁄ for estimating the carrier effective masses.

133

Expt.
Calc.
PBEsol-D3 PBEsol-D3 + U
a [Å] 22.52 23.133 23.135
c [Å] 3.3 3.140 3.438
c/a 0.147 0.136 0.149
V [Å
3
] 1449 1455 1593
M [BM] 1.87
2.15
(0.66, 0.67, 0.66)
4.77
(1.81, 1.81, 1.81)
Table 4.7. Calculated lattice parameters and magnetic moments of the CoTHT analogue of the
FeTHT framework obtained using the PBEsol-D3 functional with and without a Hubbard
correction of U = 5 eV applied to the Co d states. For each calculation, both the total magnetic
moment and the moments of the three Co ions are given. Experimental values are shown in the
second column for comparison.

Figure 4.38. Band dispersion and electronic density of states (DoS) curves of the Co analogue of
the FeTHT framework calculated with PBEsol (top) and PBEsol+U with a Hubbard correction of
134

U = 5 eV applied to the Fe d states (bottom). As in Figure 4.31, the blue and red colours denote
electronic states in the two spin channels, and the thick black lines indicate parts of the dispersion
used to evaluate 𝜕 2
𝐸 (𝑘 ) 𝜕 𝑘 2
⁄ for estimating the carrier effective masses.

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CHAPTER 5

Conductivity in a series of Metal–Benzenehexathiolate Two-Dimensional
Coordination Polymers

141

5.1 Introduction

Metal-organic frameworks, or MOFs, are a class of coordination polymers exhibiting regular,
crystalline networks with low defect concentrations, permanent porosity, and high internal surface
areas.
1-4
The first reports of MOFs occurred over 20 years ago
5
as a special example of a man-
made zeolite structures and have been investigated for their potential applications in gas storage,
separation, and catalysis.
6-10
An attractive feature of MOFs is their synthetic tunability and ease of
functionalization. MOFs can be tuned by altering either the metal centers/nodes or the organic
linker, or they can have their structures modularly tuned using principles of reticular chemistry.
1-
4,10,11
Most MOFs are formed through solvothermal methods, which encourage reversible bond
formation between the carboxylate oxygen atoms of a linker with the metal center. However, the
weak covalent overlap between the orbitals of the carboxylate oxygen atoms and those of the metal
centers leads to poor conductivity and low charge carrier mobilities, exhibited by flat, low-
dispersion bands in their electronic band structures, resulting in frameworks which act as
insulators.
11-13
Conductive MOFs could be used in a variety of device applications, such as
chemical sensing, electronics, and electrocatalysis; however, their intrinsically low conductivity
and insulating nature have greatly limited these applications.
13,14

Recently, several developments in the synthesis of new MOFs have resulted in frameworks with
appreciable electrical conductivity. Strategies for extra-framework modifications which yield
conductivity enhancements include guest species/doping,
15-21
growth of conductive polymers
through the MOF pores,
22
and the generation of composites using conductive carbon additives like
graphene.
23,24
However, these strategies can block access to the pores or greatly reduce the
accessible surface area of the MOF; as a result, the development of MOFs which are intrinsically
conductive is a desirable goal.
13,14

The use of redox-active linkers has led to massive improvements in intrinsically conductive MOFs.
Three-dimensional frameworks have shown that softer, more covalent ligands and mixed-valency
at the metal sites can give rise to 3D structures with permanent porosity and appreciable
conductivity.
18,25,26
The reduction in dimensionality to planar, two-dimensional frameworks using
semiquinones/cathecolates,
27-31
diimines,
32-37
and dithiolenes,
38-52
has been shown to dramatically
142

increase conductivity values. Benzenehexathiolate-based 2D MOFs have been shown to act as
particularly conductive frameworks; NiBHT frameworks exhibit high conductivity that can be
adjusted by chemical reduction and oxidation
40,41
, and CuBHT forms dense 2D sheets which have
conductivity values as high as 2,500 S/cm.
38,39
This is record among MOFs, and the metallic
character observed in the CuBHT framework is only the second example of metallic character (the
first being our work on the CoTHT framework
50
) and the first example of a MOF with a transition
to a superconducting regime at temperatures below 1 K.
39

Because nickel, palladium, copper, and silver have been combined with BHT to form some of the
most conductive 2D coordination polymers currently known, we were motivated to synthesize and
study the analogous iron and cobalt 2D frameworks. In light of our recent work on the metallic
conductivity observed in the CoTHT and FeTHT frameworks, we sought to understand whether
the BHT-series of 2D MOFs was capable of producing metallic conductivity and if there were any
observable trends along the 1
st
row of the transition metals. This chapter explores the conductivity
of the iron, cobalt, and nickel benzenehexathiolate 2D coordination polymers (Scheme 5.1).

Scheme 5.1. Proposed structure of the 2D metal dithiolene frameworks [M3(BHT)2]
3-
studied
here

143

5.2 Results and Discussion

The 2D MOFs in this work FeBHT, CoBHT, and NiBHT were synthesized using a liquid-liquid
interfacial method adapted from Nishihara.
39-41,49
The resultant black films were either deposited
on glass slides via direct deposition or collected as powders by decanting the liquids and collecting
the powders by filtration. Structural characterization was first attempted using powder diffraction,
but the observed patterns exhibited no peaks, indicating a lack of long-range order. Nishihara
reports weak peaks for the NiBHT coordination nanosheet, but our attempts to reproduce this on
the FeBHT or CoBHT did not yield similar results. Despite the lack of long-range order by XRD,
SEM images reveal flat, sheet-like morphologies for all three materials which are consistent with
previous reports of these MOFs (Figures 5.1-5.3).

Figure 5.1. Representative SEM images of FeBHT.

Figure 5.2. Representative SEM images of CoBHT
144

Figure 5.3. Representative SEM images of NiBHT

XPS data for films of the 2D MOFs exhibit the expected peaks for carbon, sulfur, and the metal
used to generate the framework (Figures 5.4-5.6). The three compositions were examined using
XPS. Survey spectra for FeBHT revealed the presence of Fe, S, and C. Quantification of the Fe2p
and S2s regions revealed that the Fe:S ratio was approximately 1:4, as is expected. In the iron region
(Figure 5.4b), deconvolution of the Fe 2p3/2 peak revealed three peaks at 707.5 eV, 709.2 eV, and
712.1 eV. This can be best explained as a mixture of Fe
2+
/Fe
3+
, where the peak at 707.5 eV results
from Fe
2+
, Fe
3+
is present at 709.2 eV, and the peak at 712.1 eV is a hybrid satellite primarily from
the Fe
2+
. Quantification of these peaks suggests that the ratio of Fe
2+
/Fe
3+
is approximately 2:1.
Similar behavior is observed for the cobalt region in CoBHT (Figure 5.5b): Three peaks can be fit
to the Co2p3/2 region at 779.0, 781.0, 787.2 eV . The main peak at 779.0 eV is typical of cobalt in a
+2 or +3 oxidation state; the weak satellites at 781.0 and 787.2 eV are also consistent with this
assignment. Lastly, the Ni 2p region for the NiBHT coordination polymer exhibits features for the
Ni2p3/2 at 853.0, 854.8, 860.6 eV (Figure 5.6b). The latter two peaks at higher binding energies are
satellite peaks consistent with both Ni
2+
and Ni
3+
.

145

Figure 5.4. XPS spectra for FeBHT. (a) survey scan, (b) Fe 2p, (c) S 2s, (d) S 2p

Figure 5.5. XPS spectra for CoBHT. (a) survey scan, (b) Co 2p, (c) S 2s, (d) S 2p

146

Figure 5.6. XPS spectra for NiBHT. (a) survey scan, (b) Co 2p, (c) S 2s, (d) S 2p

Powders of FeBHT, CoBHT, and NiBHT were subjected to temperature-dependent susceptibility
measurements and each shows a paramagnetic response characteristic of localized moments
(Figure 5.7-5.9). The susceptibility data show that FeBHT (and to a lesser extent CoBHT) exhibit
ordering behavior characteristic of spin glasses, whereas NiBHT exhibits no ordering. Hysteresis
loops collected at 2 K show a reduction in hysteresis across the series, with FeBHT exhibiting a
well-defined hysteresis and NiBHT exhibiting no hysteresis. Fits of the high temperature magnetic
susceptibility data (200− 300 K) for each of these samples to the Curie−Weiss equation (1)(χ =
C/(T − θCW) + χ0) show that while all three materials behave as paramagnets, FeBHT and CoBHT
have very negative Curie-Weiss temperatures (FeBHT = -323.41 K, CoBHT = -111.42 K)
demonstrating antiferromagnetism, NiBHT has a positive Curie-Weiss temperature of 34.54 K,
consistent with ferromagnetism. In the case of FeBHT, the experimental moment is consistent
with the assignment of mixed oxidation states, with both Fe
2+
and Fe
3+
present in a roughly 7:1
ratio (88%:12%). Additionally, the hexagonal network of these materials leads to Kagome nets,
which can lead to magnetic frustration. For FeBHT, the very negative θCW of -323.41 K and
ordering temperature of 24 K lead to a large frustration parameter of 13.5. For the CoBHT, the
147

experimental moment is also consistent with mixed oxidation states, suggesting that both Co
2+
and
Co
3+
are present. And for the NiBHT, the ratio of Ni
2+
to Ni
3+
is about a 3:1.

Figure 5.7. Magnetic studies of FeBHT. (left) temperature-dependent susceptibility, (center)
Curie-Weiss fit of the high-temperature region (200–300 K), and (right) hysteresis loop collected
at 2 K.

Figure 5.8. Magnetic studies of CoBHT. (left) temperature-dependent susceptibility, (center)
Curie-Weiss fit of the high-temperature region (200–300 K), and (right) hysteresis loop collected
at 2 K.

148

Figure 5.9. Magnetic studies of NiBHT. (left) temperature-dependent susceptibility, (center)
Curie-Weiss fit of the high-temperature region (200–300 K), and (right) hysteresis loop collected
at 2 K.

The temperature-dependent resistivity of each material was measured using a four-point Van der
Pauw geometry on several thin films with varying thicknesses. Graphitic carbon paint (Alfa Aesar)
was used to create Ohmic contacts, as verified by the linear temperature-dependent I-V traces
shown in Figures 5.13, 5.17, and 5.20. For films of FeBHT, conductivity values at 300 K for
different film thicknesses range from 0.06 to 13.8 S/cm, suggesting that the Fe-containing
nanosheets are less conductive than the reported Ni nanosheets (see Table 5.1). Their thicknesses
were confirmed using atomic force microscopy (Figure 5.12). An exponential rise in resistivity is
observed at high temperatures, which is consistent with semiconducting behavior (Figure 5.10).
The energy of activation for conduction in this film was 56.5 meV. Upon further cooling, all
FeBHT films (barring one 200 nm sample where electrical contact was lost at relatively high
temperature) exhibited semiconductor-metal transitions (Figure 5.11). We have previously
reported such behavior in an analogous material, a Co3(2,3,6,7,10,11-triphenylenehexathiolate)2
two-dimensional metal-organic framework (MOF), and attributed this transition to an eclipsed
stacking configuration which allows for charge transport along the c-axis.
50
Interestingly, the
transition temperature Tmetallic decreases from 220 K to 105 K with a corresponding increase in
film thickness from 25 nm to 89 nm (Table 5.1). This suggests that the film thickness plays a direct
role in the charge transport, possibly by correlating with the number of defects in the framework.

149

Figure 5.10. Temperature-dependent resistivity of a 200 nm thick sample of FeBHT (inset:
Arrhenius plot for FeBHT; the calculated Ea for this sample is 56.5 meV)

Figure 5.11. Temperature-dependent resistivity of a 25 nm thick sample of FeBHT (inset:
Arrhenius plot for FeBHT; the calculated Ea for this sample is 83.2 meV)

Figure 5.12. Atomic Force Microscopy (AFM) images of an FeTHT film with a thickness of 25
nm.

150

Figure 5.13. I-V Traces at different temperatures for a 200 nm thick film of FeBHT

Films of CoBHT ranging in thicknesses between 26.8 and 400 nm at 300 K exhibit conductivity
between 9.1 and 47 S/cm, making CoBHT more conductive than FeBHT and less conductive than
NiBHT. Their thicknesses were confirmed using atomic force microscopy (Figure 5.16).
Temperature-dependent resistivity for CoBHT also exhibits semiconducting behavior at higher
temperatures, and for relatively thick samples no transition is observed even at low T (~25 K,
Figure 5.14). Interestingly, as the CoBHT film thickness is reduced, a semiconducting-metallic
transition occurs at temperatures below 100 K (Figure 5.15). Also interesting is that the transition
temperature Tmetallic increases with decreasing film thickness, suggesting that the film thickness
plays a direct role in the charge transport, possibly by correlating with the number of defects in the
framework, as with the FeBHT. For these samples, the activation energy for charge hopping was
determined by fitting the semiconducting region to the Arrhenius expression; the results show that
the activation energy Ea decreases with increasing film thickness, with values ranging from 42.9
meV in the thinnest measured sample (26.8 nm, Figure 5.15) to 13.2 meV in the thickest sample
(400 nm, Figure 5.14).

151

Figure 5.14. (left) Temperature-dependent resistivity of a 400 nm thick sample of CoBHT.
(right) Arrhenius plot for CoBHT; the calculated Ea for this sample is 13.2 meV.

Figure 5.15. (left)Temperature-dependent resistivity of a 26.8 nm thick sample of CoBHT.
(right) Arrhenius plot for CoBHT; the calculated Ea for this sample is 42.9 meV.

Figure 5.16. Atomic Force Microscopy (AFM) images of a CoTHT film with a thickness of 94.5
nm.

152

Figure 5.17. I-V Traces at different temperatures for a 400 nm thick film of CoBHT

The high conductivity of NiBHT has been reported previously
40,41
. However, it is worth noting
that even for thin (22.2 nm) films of NiBHT, a semiconducting-metallic transition was not
observed in any sample (Figure 5.18). Atomic force microscopy images of a NiTHT film with a
thickness of 22.2 nm are shown in Figure 5.19, and a series of NiBHT film thicknesses and
conductivity values is given in table 5.1. A comparison of the Tmetallic onset for the thin samples of
FeBHT, CoBHT, and NiBHT is given in Figure 5.21. XPS of each of the three frameworks after
conductivity studies shows only minor changes (Figures 5.22-5.24).

Figure 5.18. (left) Temperature-dependent resistivity of a 22.2 nm thick sample of NiBHT.
(right) Arrhenius plot for NiBHT; the calculated Ea for this sample is 10.5 meV.

153

Figure 5.19. Atomic Force Microscopy (AFM) images of a NiTHT film with a thickness of 22.2
nm.

Figure 5.20. I-V Traces at different temperatures for a 190 nm thick film of NiBHT

Figure 5.21. Comparison of Tmetallic for FeBHT (red, 25 nm, 220 K), CoBHT (blue, 26.8 nm, 60
K), and NiBHT (green, 22.2 nm, no Tmetallic observed). Y-axis is scaled resistivity (arb. units)
154

Figure 5.22. XPS spectra for FeBHT after conductivity studies. (a) Fe 2p, (b) S 2s, (c) S 2p

Figure 5.23. XPS spectra for CoBHT. (a) Co 2p, (b) S 2s, (c) S 2p

155

Figure 5.24. XPS spectra for NiBHT. (a) Ni 2p, (b) S 2s, (c) S 2p

These experimental results should be examined in comparison with various density functional
theory (DFT) studies of the metal dithiolene MOFs reported here. Nishihara and coworkers
reported DFT studies alongside their experimental results on NiBHT which suggested that it
should exhibit metallic behavior; however, their experimental results, which we have reproduced,
show that the NiBHT framework is a narrow bandgap semiconductor.
40,41
In contrast, another DFT
study on the NiBHT system suggests that it should exhibit semiconducting behavior.
53
DFT has
also been reported on the FeBHT and CoBHT frameworks; however, their calculated magnetic
behavior (ferromagnetic) contrasts with our experimental data (antiferromagnetic).
54
Further
experimental and computational studies of these two-dimensional systems is desirable.

156

Compound Thickness (nm) ρ (300 K, Ω·cm) σ (300 K, S/cm) Tmetallic (K) Ea (meV)
FeBHT
700 1.6 0.625 105 140
200 6.38 0.16 - 56.5
89 0.145 6.88 90 104
45 15.66 0.0638 180 144
25 0.08 13.8 220 83.2
CoBHT
400 0.0213 47 - 13.2
200 0.0375 26.7 20 14.8
94.5 0.081 13.0 25 21.6
26.8 0.11 9.1 60 42.9
NiBHT
190 0.00707 141 - 18.7
83 0.026 27.8 - 88.4
22.2 0.00386 259 - 10.5
Table 5.1. Summary of electrical transport behavior for FeBHT, CoBHT, and NiBHT

5.3 CONCLUSIONS

In summary, we have synthesized, characterized, and investigated the temperature-dependent
resistivity for a series of metal (M = Fe, Co, Ni) benzenehexathiolate frameworks. The Tmetallic
transition temperature appears to across the row. Sample to sample variance prevents more
sophisticated claims, such as the thickness-dependence of the conductivity. These results identify
experimentally observed MOFs that exhibits band-like metallic conductivity, and suggest that
choice of metal in these systems can allow for tuning of metallic behavior. We expect the design
principles discovered in these fundamental studies to have an impact on our understanding of the
charge transport characteristics of MOFs, leading to new materials with impressive electrical
properties and applications in electronic devices.

5.4 EXPERIMENTAL DETAILS

5.4.1 General Considerations

All manipulations of air and moisture sensitive materials were conducted under a nitrogen
atmosphere in a Vacuum Atmospheres drybox or on a dual manifold Schlenk line. The glassware
was oven-dried prior to use. Dichloromethane was degassed with nitrogen and passed through
activated alumina columns and stored over 4 Å Linde-type molecular sieves. Ethyl acetate, water,
157

and methanol were placed under vacuum and refilled with nitrogen (10 ⨯). Benzenehexathiol
(BHT) was prepared according to the reported procedures. Water was deionized with the Millipore
Milli-Q Synergy system (18.2 MΩ·cm resistivity). All other chemical reagents were purchased
from commercial vendors and used without further purification.

5.4.2 Synthesis of FeBHT

Fe3(benzenehexathiolate)2 (FeBHT): 5 mL of dichloromethane was added to 2.0 mg of BHT in a
20 mL scintillation vial in a nitrogen glovebox. The resulting suspension was sealed and sonicated
for 15 minutes, then added to 40 mL of dichloromethane in a 120 mL glass jar. Water (40 mL)
was gently added to the organic solution to create a liquid-liquid interface. An aqueous solution of
FeCl2·4H2O (40 mg, 0.2 mmol, 5 mM) and was gently added to the aqueous layer and the jar was
sealed and allowed to stand overnight. For thinner films, the thickness was controlled by reducing
the amount of BHT in the reaction.

5.4.2 Synthesis of CoBHT

Co3(benzenehexathiolate)2 (CoBHT): 5 mL of dichloromethane was added to 2.0 mg of BHT in a
20 mL scintillation vial in a nitrogen glovebox. The resulting suspension was sealed and sonicated
for 15 minutes, then added to 40 mL of dichloromethane in a 120 mL glass jar. Water (40 mL)
was gently added to the organic solution to create a liquid-liquid interface. An aqueous solution of
CoCl2·6H2O (50 mg, 0.21 mmol, 5.25 mM) and was gently added to the aqueous layer and the jar
was sealed and allowed to stand overnight. For thinner films, the thickness was controlled by
reducing the amount of BHT in the reaction.

5.4.2 Synthesis of NiBHT

Ni3(benzenehexathiolate)2 (NiBHT) was prepared according to reported methods.
40,41

Scanning Electron Microscopy (SEM) images were collected using a JEOL-7001F microscope
operating at 15 kV with 4 nA of probe current.

XPS data were collected using a Kratos AXIS Ultra instrument. The monochromatic X-ray source
was the Al K α line at 1486.6 eV, and the hybrid lens and slot mode were used. Low-resolution
158

survey spectra were acquired between binding energies of 1–1200 eV. Higher-resolution detailed
scans, with a resolution of ~0.2 eV, were collected on individual XPS regions of interest. The
sample chamber was maintained at < 9⨯10
–9
Torr. The XPS data were analyzed using the CasaXPS
software.

Conductivity measurements were performed using a custom set up integrated into a 14T Quantum
Design Dynacool Physical Properties Measurement System. A Keithley 6220 Precision Current
Source (excitation currents of 1-100 nA) was used to trigger and control a Keithley 2182A
nanovoltmeter. In order to minimize errors associated with contact resistance and drift voltages, a
Keithley 2172 matrix switch equipped with a Keithley 6536 Hall effect card was used to alternate
the direction of the applied current. Because of difficulty associated with preparing samples with
uniform dimensions, all measurements were performed in a four-point probe Van der Pauw
geometry. Copper wire contacts were attached to the films using a conductive carbon paint (Alfa
Aesar) and soldered onto a Quantum Design puck with resistivity option. All measurements were
performed under a reduced pressure of ~10 torr.

Atomic Force Microscopy (AFM) topography images were collected using an Agilent 5420 SPM
instrument operating in tapping mode. The probe tips were Tap300-G Silicon AFM probes
(resonant frequency 300 kHz, force constant 40 N/m) purchased from Budgetsensors.com and
aligned prior to use. Images were collected with a scan rate of 0.1 sec/line and over an area of 40
µm. All samples were imaged under one atmosphere of air at room temperature.

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Abstract (if available)

Abstract Electrocatalytic hydrogen evolution from water using solar energy is one potential pathway towards a clean energy economy. The work in this dissertation attempted to improve upon the stability and activity of a well-studied hydrogen-evolving cobalt dithiolene electrocatalyst by incorporating it into a metal-organic framework (MOF). This work shows that metal dithiolene units can be successfully integrated into two-dimensional frameworks, and that the generated coordination polymers function as efficient electrocatalysts for the hydrogen evolving reaction under fully aqueous conditions. Their electrical conductivity was also investigated and shown to be among that of the best coordination polymers. The design principles demonstrated in these studies have had a significant impact on the development of other catalytically-active frameworks in the field and will hopefully lead to more efficient devices for solar energy conversion and storage.

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Asset Metadata

Creator Clough, Andrew James (author)

Core Title Two-dimensional metal dithiolene metal-organic frameworks as conductive materials for solar energy conversion

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 07/24/2019

Defense Date 06/11/2019

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

Tag 2D MOF,conductivity,conversion,dithiolene,energy,HER,hydrogen evolution,metallic,metal-organic framework,MOF,OAI-PMH Harvest,solar,two-dimensional

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Marinescu, Smaranda (committee chair), Melot, Brent (committee member), Ravichandran, Jayakanth (committee member)

Creator Email ajclough@usc.edu,clough.andrew.james@gmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c89-187932

Unique identifier UC11662461

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Legacy Identifier etd-CloughAndr-7599.pdf

Dmrecord 187932

Document Type Dissertation

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Rights Clough, Andrew James

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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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