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Solution phase synthesis routes to functional nanomaterials for energy storage

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Solution phase synthesis routes to functional nanomaterials for energy storage

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SOLUTION PHASE SYNTHESIS ROUTES TO FUNCTIONAL
NANOMATERIALS FOR ENERGY STORAGE

by
Gözde Barim
__________________________________________________________________
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)

December 2018

Copyright 2018 Gözde Barim

ii

To my irrational and lonely dreamland
Şuursuzya/Irrationaland

iii

Acknowledgements

This dissertation would have not been possible without the support of many benevolent
fellows. First and foremost, I would like to express my profound gratitude to my thesis
advisor Professor Richard L. Brutchey. His consistent guidance helped me to enhance my
knowledge and expertise that today’s exceedingly progressive and competitive research
and work environment demand. He has created an ideal environment for me to grow both
academically and personally as well as to develop the right mind-set in order to be an
accomplished researcher. I would like to thank Richard for all that he has done for me. I
will forever be grateful to him for giving me the opportunity to work with him and for the
knowledge and skills that I have gained working under his supervision.
I also would like to sincerely thank my dissertation committee members, Professors,
Smaranda Marinescu, Sarah Feakins, Surya Prakash and Ralf Haiges for their time and
providing guidance. I am also deeply thankful for my collaborators, Professors Brent Melot
and Benjamin Morgan as well as Dr. Shiliang Zhou. I have learned a lot about energy
storage devices by collaborating with them. I am truly grateful for the time that we have
been able to work together, and the opportunity to expand my skill set and knowledge.
Besides having a terrific supervisor, collaborators and mentors, I have also been
fortunate to have awesome fellows and friends. I would like to genuinely thank all present
and former Brutchey Group members, especially, Dr. Matthew Greaney, Dr. Jannise
Buckley, Dr. Sean Culver, Dr. Haipeng Lu, Dr. Carrie McCarthy, Dr. Patrick Cottingham,
Dr. Lucia Mora, Emily Roberts, Bryce Tappan and Sara Smock for their invaluable help,
contribution and friendship. It has been a great pleasure and so much fun spending time

iv

with them and learning together. Special thanks to my dear friends from USC for their
candid friendship, moral support and being there, in every sense of the phrase. This list
includes but not limited to Damir Popov, Mu’azzam Idris, Carolina Amador, Beril
Kırağası, Nil Şimşek, Samet Keserci, Şeyma Ekiz, Şebnem Baran, Güher Çamlıyurt,
Günce Cinay, Alpi, Canan İpek and Betül Mulugün.
Finally, I would like to thank my mother and father, Hülya and Rebi Barım for enduring
love, patience and encouragement. I honestly can’t thank enough for all that my family and
friends have done for me over the last five years.

v

Table of Contents

Acknowledgements iii

List of Tables viii

List of Figures x

Abstract xvi

Chapter 1. Introduction 1
1.1. Cathode materials for rechargeable ion batteries 1
1.2. Structural and electrochemical properties of NASICON-type cathode materials 2
1.3. Binary nickel sulfide system for energy applications 8
1.4. Solution-phase synthesis of binary nickel sulfides 11
1.5. References 17

Chapter 2. Influence of Rotational Distortions on Li
+
- and Na
+
-Intercalation in Anti-
NASICON Fe2(MoO4)3 24
2.1. Abstract 24
2.2. Introduction 24
2.3. Results and Discussion 27
2.3.1. Li
+
insertion 32
2.3.2. Na
+
insertion 42
2.4. Experimental 61
2.4.1. Synthesis of Fe2(MoO4)3 61
2.4.2. Chemical insertion 61
2.4.3. Characterization 62
2.4.4. Electrochemical measurements 63
2.4.5. Symmetry-mode analysis 64
2.4.6. Computational 64
2.5. Conclusions 65
2.6. References 66

vi

Chapter 3. Investigating the Mechanism of Reversible Lithium Insertion into Anti-
NASICON Fe2(WO4)3 70
3.1. Abstract 70
3.2. Introduction 70
3.3. Results and Discussion 73
3.4. Experimental 90
3.4.1. Synthesis 90
3.4.2. Chemical Insertion 91
3.4.3. Characterization 91
3.4.4. Electrochemical Measurements 93
3.5. Conclusions 94
3.6. References 95

Chapter 4. Phase Control in the Colloidal Synthesis of Well-Defined Nickel Sulfide
Nanocrystals 99
4.1. Abstract 99
4.2. Introduction 99
4.3. Results and Discussion 102
4.3.1. Sulfur-deficient Ni-S Phases 103
4.3.2. Sulfur-rich Ni-S Phases 117
4.4. Experimental 124
4.4.1. General Considerations 124
4.4.2. Ni3S2 Nanocrystal Synthesis 124
4.4.3. Ni9S8 Nanocrystal Synthesis 124
4.4.4. Ni3S4 Nanocrystal Synthesis 125
4.4.5. α-NiS Nanocrystal Synthesis 126
4.4.6. β-NiS Nanocrystal Synthesis 126
4.4.7. Nanocrystal Purification 127
4.4.8. Instrumentation 127
4.5. Conclusions 128
4.6. References 129

vii

Chapter 5. Synthesis and Characterization of Thiospinel CoNi2S4 and FeNi2S4
Nanocrystals 133
5.1. Introduction 133
5.2. Experimental Details 135
5.2.1. General Considerations 135
5.2.2. Nanocrystal Syntheses 135
5.2.3. Characterization of Nanocrystals 136
5.3. Results and Discussion 137
5.3.1. CoNi2S4 Nanocrystals 137
5.3.2. FeNi2S4 Nanocrystals 140
5.4. Conclusions 142
5.5. References 143

Bibliography 145

viii

List of Tables

Table 2.1: Resulting unit cell, atomic positions and atomic displacement
parameters from the Rietveld refinement of Fe2(MoO4)3 against
synchrotron X-ray powder diffraction data and neutron diffraction
data (the Rwp was determined to be 6.01% for all banks being refined
simultaneously). Both sets of data were collected at RT. 29

Table 2.2: Resulting unit cell, atomic positions, and atomic displacement
parameters from the Rietveld refinement of lithiated Li2Fe2(MoO4)3
against neutron diffraction data (the Rwp was determined to be 3.02%
for all banks being refined simultaneously). 37

Table 2.3: Transformation matrix for converting parent structure (Pbcn) into
reference structure (P21/c) for symmetry mode analysis. 38

Table 2.4: Unit cell parameters for parent structure in reference cell settings. 38

Table 2.5: Atomic pairing for the parent phase and the distorted phase in the
reference settings. 39

Table 2.6: Displacement vectors for pristine Fe2(MoO4)3. 40

Table 2.7: Lattice parameters of NaxFe2(MoO4)3 from Le Bail fit of in situ XRD
patterns. x values are calculated based on number of electrons flown
into the cell. 45

Table 2.8: Unit cell parameters and atomic positions of optimized Li-inserted
Fe2(MoO4)3 from PBEsol+U DFT calculations. 47

Table 2.9: Unit cell parameters and atomic positions of optimized Na
+
-inserted
Fe2(MoO4)3 from PBEsol+U DFT calculations. 47

Table 2.10: Resulting unit cell parameters and atomic positions of the Fe2(MoO4)3
framework from Ritveld refinement agasint the neutron diffraction
data of Na
+
-inserted Fe2(MoO4)3 using calculated Na2Fe2(MoO4)3
with Na
+
ions omitted. 53

Table 2.11: Dihedral angles of O-Mo-Fe-O in refined frameworks of the Li
+
- and
Na
+
-inserted Fe2(MoO4)3. Dihedral angle values based on the Fe and
Mo on the top-left corner in Figure 2.18 are included. 57

Table 3.1: Resulting unit cell, atomic positions, and atomic displacement
parameters from the fit of the monoclinic model to the experimental
PDF of pristine Fe2(WO4)3 neutron total scattering data obtained at the
NOMAD beamline at Oak Ridge National Laboratory. 76

ix

Table 3.2: Resulting unit cell, atomic positions, and atomic displacement
parameters from the fit of the orthorhombic model to the experimental
PDF of lithiated Li2Fe2(WO4)3 neutron total scattering data obtained
at the NOMAD beamline at Oak Ridge National Laboratory. 82

Table 3.3: Transformation matrix for converting parent structure (Pbcn) into
reference structure (P21/c) for symmetry mode analysis. 86

Table 3.4: Unit cell parameters of parent structure into reference structure
settings. 86

Table 3.5: Atomic pairing for the parent phase and the distorted phase in the
reference settings. 87

Table 3.6: Displacement vectors for pristine Fe2(WO4)3 89

Table 4.1: Summary of the reaction conditions to obtain various Ni–S
nanocrystal phases 104

Table 4.2: Synthetic conditions for the preparation of shape-controlled Ni9S8
nanoparticles. NiI2 is reacted with various sulfur sources in
oleylamine at 180 ˚C. 115

x

List of Figures

Figure 1.1: NASICON (generally rhombohedral) and anti-NASICON (generally
monoclinic) frameworks of general formula AxMM′(XO4)3. Reprinted
with permission from ref. 7 (Copyright 2013 American Chemical
Society). 3

Figure 1.2: The (010) view of the structure of fully sodiated Na2Fe2(MoO4)3.
Reprinted with permission from ref. 21 (Copyright 2018 American
Chemical Society). 6

Figure 1.3: Comparison between sodiation and lithiation process in Fe2(MoO4)3.
Schematic diagrams of ‘‘discrete occupation’’ and
‘‘pseudocontinuous occupation’’ during Li and Na ions intercalation
into Fe2(MoO4)3. Solid red circles and dash green ellipses stand for
Li1 (or Na1) and Fe2(MoO4)3 frameworks, respectively. Reprinted
with permission from ref. 22 (Copyright 2015 Springer Nature
Limited). 7

Figure 1.4: Calculated condensed nickel-sulfur phase diagram by Waldner and
Pelton. Reprinted with permission from ref. 23 (Copyright 2004
Deutsche Gesellschaft für Materialkunde). 9

Figure 1.5: Figure 1.5. Unit cell structures, scanning electron microscopy images,
and XRD patterns of (a) NiS, (b) NiS2, and (c) Ni3S2. Reprinted with
permission from ref. 54 (Copyright 2016 Royal Society of Chemistry).
11

Figure 1.6: (a) TEM image of the Ni3S4 nanoprism synthesized using NiCl2 as the
Ni precursor; (b) TEM image of the Ni3S4 tetrahedron (nanopyramids)
synthesized using Ni(acac)2 as the Ni precursor; (c and d) length (c)
and width (d) distribution for the Ni3S4 nanoprisms (40 nm × 10 nm);
(e) length distribution of the edges for the Ni3S4 tetrahedra (16 nm).
Reprinted with permission from ref. 76 (Copyright 2014 Royal
Society of Chemistry). 13

Figure 1.7: Diagrammatic representation of the compositions obtained with 50
mL of 0.25 M NiCl2•6H2O (12.5 mmol) and 25-300 mL of 1 M
sodium dithionite (25-300 mmol) at various pH values. The
compositions within the rectangular boxes are crystalline Ni3Sx while
those in the shaded area are crystalline sulfur (dark area) and
amorphous or poorly crystalline NiySx (light area). Reprinted with
permission from ref. 82 (Copyright 2001 American Chemical Society)
15

xi

Figure 2.1: Illustration of (a) the unit cell of pristine, monoclinic Fe2(MoO4)3 and
(b) a “lantern unit” that consists of three MoO4 tetrahedra connecting
two FeO6 octahedra. 26

Figure 2.2: TEM image of pristine Fe2(MoO4)3. 28

Figure 2.3: Rietveld refinements of pristine, monoclinic Fe2(MoO4)3 against (a)
synchrotron XRD pattern obtained at the 11-BM beamline at Argonne
National Laboratory and (b) neutron diffraction pattern obtained at the
POWGEN beamline at Oak Ridge National Laboratory. The weighted
profile R-factor (R wp) was determined to be 6.01% for all banks being
refined simultaneously and weighed equally. Absorption correction
was carried out using absorption function 0 in GSAS. Only the
diffraction pattern from bank 3 at POWGEN is given here. 29

Figure 2.4: Raman spectra of (a) pristine Fe2(MoO4)3 under ambient conditions
and (b) Li
+
-inserted and Na
+
-inserted Fe2(MoO4)3 along with the
pristine Fe2(MoO4)3 in the air-free quartz cell. *Indicates band
corresponding to air-free cell. 31

Figure 2.5: High-resolution XPS spectra of pristine Fe2(MoO4)3. 32

Figure 2.6: (a) Galvanostatic electrochemical cycling of Fe2(MoO4)3 against Li
+

insertion and its derivative (shown as inset). (b) 2D pattern based on
the in situ XRD of Li
+
insertion into Fe2(MoO4)3. 33

Figure 2.7: Rietveld refinement of one-hour X-ray diffraction pattern on pristine
Fe2(MoO4)3. 34

Figure 2.8: Rietveld refinement of one-hour X-ray diffraction pattern on fully
electrochemically lithiated Li2Fe2(MoO4)3. Intensities offset at 30, 45,
47, and 49 degrees are contributed by the beryllium cell used for in
situ X-ray diffraction. 35

Figure 2.9: Crystal structure of (a) pristine, monoclinic Fe2(MoO4)3 and (b) fully
lithiated, orthorhombic Li2Fe2(MoO4)3. 36

Figure 2.10: Refinement results of neutron powder diffraction patterns of
chemically lithiated Li2Fe2(MoO4)3. 36

Figure 2.11: Amplimodes analysis between pristine Fe2(MoO4)3 and lithiated
Li2Fe2(MoO4)3, with positions of Fe, Mo, and O in pristine
Fe2(MoO4)3 and position of Li in Li2Fe2(MoO4)3 converted into the
same reference structure setting. The transformation vectors are
plotted on each atom in (a) and filtered by amplitude of 0.4 Å in (b)
for clarity. FeO6 octahedra are shown in green, and MoO4 tetrahedra

xii

are shown in gray. Li
+
ions are left out for amplimodes analysis but
are drawn in the reference structure for clarity. 41

Figure 2.12: (a) Galvanostatic electrochemical cycling of Fe2(MoO4)3 against Na
+

insertion and its derivative (shown as inset). There are two slope
regions during both insertion and de-insertion. The turn from one
region to the other corresponds to 0.8 Na
+
per formula being
(de)inserted. (b) 2D pattern based on the in situ XRD of Na
+
insertion
into Fe2(MoO4)3. 43

Figure 2.13: Lattice evolution along Na
+
insertion into Fe2(MoO4)3 from Le Bail
fitting of in situ X-ray diffraction patterns. 44
Figure 2.14: Rietveld refinement against neutron diffraction data of the chemically
Na
+
-inserted Fe2(MoO4)3 collected at the POWGEN beamline at the
Oak Ridge National Laboratory using the calculated Na2Fe2(MoO4)3
model listed in Table S9 (χ
2
= 6.393, Rwp = 4.43%). 51

Figure 2.15: Rietveld refinement of the Fe2(MoO4)3 framework derived from the
calculated Na2Fe2(MoO4)3 (χ
2
= 6.399, Rwp = 4.43%). 52

Figure 2.16: Resulting Fe2(MoO4)3 framework from Rietveld refinement against
neutron diffraction data using optimized Na2Fe2(MoO4)3 from
simulations with sodium omitted as the starting model. 52

Figure 2.17: Experimental total scattering data for pristine, Li
+
-inserted, and Na
+
-
inserted Fe2(MoO4)3. 54

Figure 2.18: Illustration of a representative unit of the refined framework of Li
+
-
and Na
+
-inserted Fe2(MoO4)3. Li
+
from Rietveld refinement of
Li2Fe2(MoO4)3 is illustrated for perspective. 56

Figure 3.1: XRD patterns of as-prepared powders annealed at various
temperatures. 74

Figure 3.2: (a) Fit of the monoclinic model to the experimental PDF of pristine
Fe2(WO4)3 neutron total scattering data obtained at the NOMAD
beamline at Oak Ridge National Laboratory and (b) the resulting
crystal structure. (Iron, tungsten and oxygen are shown as green,
purple and orange, respectively.) 75

Figure 3.3: High-resolution XPS spectra of Fe (left) and W (right) in pristine
Fe2(WO4)3. 77

Figure 3.4: TEM images of Fe2(WO4)3 (a) before and (b) after annealing to 550
˚C. 78

xiii

Figure 3.5: (a) Galvanostatic electrochemical cycling of Fe2(WO4)3 against Li
+

insertion and (b) its derivative curve. 79

Figure 3.6: 2D pattern of in situ XRD study for a full discharge/charge cycle of
Fe2(WO4)3 against lithium at C/10 rate with a CoKα source (λ1 =
1.78897 Å, λ2 = 1.79285 Å). Color indicates the intensities of
reflections. 81

Figure 3.7: (a) Fit of the orthorhombic model to the experimental PDF of lithium-
inserted Li2Fe2(WO4)3 neutron total scattering data obtained at the
NOMAD beamline at Oak Ridge National Laboratory and (b)
resulting crystal structure. (Iron, tungsten, oxygen and lithium are
shown as green, purple, orange and yellow, respectively.) 83

Figure 3.8: Normalized (a) Fe K-edge and (b) W LIII-edge XANES spectra, and
(c) experimental neutron total scattering data for Fe2(WO4)3 and
chemically lithiated Li2Fe2(WO4)3. 84

Figure 3.9: The transformation vectors on each atom from the symmetry-mode
analysis between pristine Fe2(WO4)3 and lithiated Li2Fe2(WO4) as
distorted phase and parent phase, respectively. Positions of Fe, W, and
O atoms in Fe2(WO4)3 and position of lithium in Li2Fe2(WO4)3
converted into the same reference structure setting (P21/c). Li
+
ions
are left out for the symmetry-mode analysis but are drawn in the
reference structure for clarity. (Iron, tungsten, oxygen, and lithium
atoms are shown as green, purple, orange, and yellow, respectively.) 85

Figure 4.1: Reaction pathways indicating the preparation of various Ni-S
nanocrystals in the presence or absence of 1-dodecanethiol as the
reactivity-directing agent. 103

Figure 4.2: (a) XRD pattern of rhombohedral Ni3S2 nanocrystals. (b) TEM
micrograph of multipod-type Ni3S2 nanocrystals. (c-d) HR-TEM
micrographs of an individual Ni3S2 nanocrystal. 104

Figure 4.3: (a) Powder XRD patterns and (b) TEM micrographs of the products
from the reaction of NiI2 and N,N’-dibutyl thiourea (BuThU) in the
presence of various amounts of 1-dodecanethiol (DDT) at 180 ˚C for
1 h. 105

Figure 4.4: (a) XRD pattern of orthorhombic Ni9S8 nanocrystals. (b) TEM
micrograph of 8.8-nm Ni9S8 nanocrystals. (c) HR-TEM micrograph
of a single Ni9S8 nanocrystal. (d) Size histogram showing the
distribution of particle diameters for Ni9S8 (N = 355). 107

xiv

Figure 4.5: (a) Powder XRD patterns and (b) TEM micrographs of the products
from the reaction of NiI2 and 3.0 molar equivalents of N,N’-diphenyl
thiourea in the presence and absence of 1-dodecanethiol (DDT) in
oleylamine. 108

Figure 4.6: (a) Powder XRD patterns and (b) TEM micrographs of the products
from the reaction of NiI2 and 3.0 molar equivalents of N,N’-diphenyl
thiourea (DPhT) in the presence of 3.0 mL of 1-dodecanethiol (DDT)
in oleylamine. 108

Figure 4.7: XRD patterns of the products from the reaction of NiI2 and 3.0 molar
equivalents of N,N’-diphenyl thiourea at 180 ˚C in the (a) absence and
(b) presence of 1-dodecanethiol under otherwise identical conditions. 110

Figure 4.8:
1
H NMR of spectra in dichloromethane-d2 of N,N’-diphenyl thiourea
heated in oleylamine at 180 ˚C for 10 min, and then 5 min after 1-
dodecanethiol (DDT) was added. (a) and (b) are showing full spectra
and aromatic region only, respectively.
1
H NMR spectra of starting
materials (i.e., oleylamine (OAm) (yellow), N,N’-diphenyl thiourea
(PhThU) (pink), and 1-dodecanethiol (DDT) (blue)). Residual solvent
is denoted by *. 112

Figure 4.9: Powder XRD patterns of the products from the reaction of NiI2 and
various amounts of 1-dodecanethiol in oleylamine (OLA). 113

Figure 4.10: TEM micrographs of Ni9S8 nanocrystals synthesized by the reaction
of NiI2 with 1-dodecanethiol at 180 ˚C for 2 h with thiol:Ni ratios of
(a) 1.5, (b) 3.0 and (c) 6.0. Histograms of the particle length
distributions for nanocrystals synthesized with (d) 1.5, (e) 3.0, and (f)
6.0 molar equivalents of 1-dodecanethiol. 114

Figure 4.11: Powder XRD patterns of the products from the reaction of NiI2 with
variable amounts of N,N’-diphenyl thiourea: (a) as a function of sulfur
S:Ni ratio; all reactions were carried out at 180 ˚C for 4 h, and (b) as
a function of temperature for 4 h with a S:Ni molar ratio of 4.0. 119

Figure 4.12: (a) XRD pattern of cubic Ni3S4 nanocrystals. (b) TEM micrograph of
12.4-nm Ni3S4 nanocrystals. (c) HR-TEM micrograph of an individual
Ni3S4 nanocrystal. (d) Size histogram showing the distribution of
nanocrystal diameters for Ni3S4 (N = 301). 120

Figure 4.13: (a) XRD pattern of hexagonal α-NiS nanocrystals. (b) TEM
micrograph of 8.9-nm α-NiS nanocrystals. (c) HR-TEM micrograph
of an individual α-NiS nanocrystal. (d) Size histogram showing the
distribution of nanocrystal diameters for α-NiS (N = 300). 122

xv

Figure 4.14: (a) Powder XRD patterns of rhombohedral β-NiS nanocrystals
synthesized by N,N’-diphenyl thiourea with the α-NiS impurity shown
by (*). (b) TEM micrograph of β-NiS nanocrystals. 123

Figure 5.1: (a) Powder XRD patterns of cubic spinel CoNi2S4 nanocrystals
synthesized by 1-dodecanethiol. (b) TEM micrograph of CoNi2S4
nanocrystals. 137

Figure 5.2: XRD patterns of the aliquots from the reaction of Co(acac)2 and
Ni(acac)2 in oleylamine and 1-dodecanethiol at 180 ˚C at various
reaction times. 138

Figure 5.3: XRD patterns of the aliquots from the reaction of Co(acac)2, Ni(acac)2
and N,N’-diphenyl thiourea in oleylamine at 180 ˚C with various
reaction times. 139

Figure 5.4: (a) Powder XRD patterns of cubic spinel FeNi2S4 nanocrystals
synthesized by 1-dodecanethiol. (b) TEM micrograph of FeNi2S4
nanocrystals. 140

Figure 5.5: (a) Powder XRD patterns of cubic spinel FeNi2S4 nanocrystals
synthesized by 1-dodecanethiol with various reaction times. (b) TEM
micrograph of FeNi2S4 nanocrystals with a 3 h reaction. 141

xvi

Abstract

Electrochemical, magnetic, and electronic properties of functional nanomaterials
depend on their size, shape, composition, and crystal phase. Therefore, having a high
degree of control over the synthesis of functional nanomaterials is significant for practical
applications. The majority of synthetic approaches for functional nanomaterials require
highly energy- and capital-intensive conditions and, as such, there is a need to develop new
methodologies for the synthesis of functional nanomaterials under more benign conditions.
In this work, large-scale, solution-based syntheses of anti-NASICON Fe2(MoO4)3 and
Fe2(WO4)3 nanoparticulates, as well as synthetic control over nanocrystal phase in the
complicated binary nickel–sulfur phase space, have been featured.
The correlation of structural changes in the frameworks of anti-NASICON type
polyanionic cathode materials with their electrochemical properties has been demonstrated
through a wide spectrum of structural and electrochemical characterization techniques.
Two-phase structural change upon lithium insertion into Fe2(MO4)3, where M = Mo or W,
has been confirmed by in situ powder X-ray experiments and galvanostatic electrochemical
cycling. Whereas, sodium insertion into Fe2(MoO4)3 proceeds trough a single-phase, solid-
solution type mechanism. To gain better insight on the mechanistic difference between
lithium and sodium insertion into Fe2(MoO4)3 a combination of high-resolution diffraction
techniques and neutron total scattering techniques have been used. Symmetry-mode
analysis demonstrated that concerted polyhedral rotations facilitates lithium and sodium
insertion into these polyanionic cathode materials. In the case of lithium insertion, the
monoclinic pristine structure transforms into a lithiated orthorhombic structure via the
polyhedral rotational distortions due the electrostatic interactions between oxygens in the

xvii

polyhedral subunits and inserted lithium ions. Although sodium insertion into Fe2(MoO4)3
is fundamentally similar, the larger ionic radius of sodium restricts concerted polyhedral
rotations and results in a less ordered sodiated structure.
Independent preparation of morphologically well-defined nanocrystals of Ni3S 4, NiS,
Ni9S 8, and Ni 3S 2 has also been presented via a systematic evaluation of factors that
simultaneously control the phase and yield well-defined nickel sulfide nanocrystals. Conversion
kinetics of N,N’-disubstituted thioureas have been found as the main factor for the nanocrystal
phase control in the binary Ni-S system. Relatively less reactive N,N’-butyl thiourea leads to
formation of the sulfur-deficient Ni9S 8 phase, whereas, more reactive N,N’-phenyl thiourea
favors stoichiometric NiS and sulfur-rich Ni3S 4 phases. Employment of 1-dodecanethiol
during the nanocrystal growth stage impede sulfur incorporation into growing nanocrystals and
most sulfur-deficient Ni3S 2 has been obtained by using N,N’-butyl thiourea in the presence of
1-dodecanethiol. In the absence of 1-dodecanethiol, a phase evolution from Ni3S 4 to NiS
phase with an increase in the temperature and sulfur to nickel precursor ratio by using N,N’-
phenyl thiourea as the primary sulfur precursor was confirmed.
This synthetic approach has then been expanded to mixed-metal thiospinels. Cubic
spinel CoNi2S4 and FeNi2S4 nanoparticles have been synthesized by the fast-injection of 1-
dodecanethiol into a hot mixture of metal precursors in oleylamine at 180 ˚C. Structural
and morphological characterization of as-synthesized nanocrystals has been demonstrated
as well.

1

Chapter 1. Introduction

1.1. Cathode materials for rechargeable ion batteries

Building a sustainable energy infrastructure is one of the most critical challenges that
we face currently. Renewable energy sources, such as solar and wind, are considered
feasible solutions to energy crisis, however, the intermittent nature of these technologies
makes it complicated to effectively utilize the energy produced by these sources due to
fluctuations in demand throughout the day. Owing to the high energy efficiencies,
scalability and modularity, rechargeable ion batteries are promising technologies for grid-
level energy storage implementations. Rechargeable ion batteries function on reversible
insertion and extraction of guest ions into/from an electrode material during discharging
and charging. In rechargeable ion batteries, diffusion of ions though an electrolyte results
in a redox reaction of the host framework assisted by a flow of electrons through an external
circuit. During discharging, guest ions are inserted into cathode leading to reduction of
metal ions in the cathode material and generating a flow of electrons from anode to cathode,
whereas charging process reverses the direction of flowing ions and electrons.
Cathodes are the key component of the rechargeable ion batteries to achieve high
capacity and superior power delivery. A majority of commercialized cathode materials in
lithium-ion batteries are based on transition metal oxides.
1,2
Oxide-based cathode materials
generally possess a layered structure and are considered the most promising cathode type
due to their favorable balancing between capacity, rate capability, voltage and density.
3

However, they suffer from structural instability and reactivity towards electrolytes.
4
In
oxide-based cathodes, the reversible insertion of guest ions is accompanied by substantial

2

changes to the unit cell that result from distortions in the structural framework causing
degradation of the performance over time.
5

As an inexpensive and safer alternative to oxide-based hosts, polyanionic cathode
materials have been extensively investigated after the discovery of olivine-type LiFePO4.
6

Polyanionic hosts comprised of three-dimensional (3D) network of transition metals and
polyanion units (XO4)
n-
that display rigid covalent bonds between main group elements and
their surrounding ligands.
7
Although, polyanionic cathodes suffer from smaller theoretical
gravimetric capacities, they possess very stable frameworks favorable for long-term
structural stability and safety.
7,8
They also exhibit versatility toward cation and anion
substitution enabling a variety of crystal structures and atomic arrangements.
9

1.2. Structural and electrochemical properties of NASICON-type cathode materials

NASICON (Na-super-ionic-conductors)-type frameworks are among the first
examples of polyanionic materials investigated as cathodes for lithium- and sodium-ion
batteries. The NASICON structure has a general formula of AxMM
’
(XO4)3 and consists of
a three-dimensional network of corner sharing octahedra (MO6 and M
’
O6) and tetrahedra
(XO4).
7
The structure of NASICON frameworks can be described by repeating units called
lantern units. Lantern units [MM
’
(XO4)3] are produced by two octahedra connected to three
tetrahedra (Figure 1.1). If lantern units stack alternately along the (2b+c) direction, the
structure is called anti-NASICON and the crystal structural usually adopts either
monoclinic (P21/a) or orthorhombic (Pcan) space groups.
7
Highly covalent 3D frameworks
constructing large interstitial voids facilitate fast ion transport in the NASICON structure.

3

Figure 1.1. NASICON (generally rhombohedral) and anti-NASICON (generally
monoclinic) frameworks of general formula AxMM′(XO4)3. Reprinted with permission
from ref. 7 (Copyright 2013 American Chemical Society).

In the 1980s, Goodenough et al. investigated iron-based NASICON type materials with
a general formula of AbFe2(XO4)3, where X = S, P, Mo, W.
10,11
These materials exhibit
higher voltages than simple oxide-based compounds upon discharging and charging, which
is attributed to covalency in the Fe-O bonds. This stems from the inductive effect of XO4
groups and strong polarization of O
2-
toward the polyanion center X
n+
, which results in
weaker Fe 3d-O 2p hybridization leading to a lower energy for Fe
2+
/Fe
3+
redox couples.
9

For instance, substitution of S for Mo decreases covalency of Fe-O bonds and lowers the
energy level of antibonding states. As a result, the energy difference between the Fe
2+
/Fe
3+

and Li
+
/Li redox couples increases leading to 0.6 V increase in the cell voltage for
Fe2(SO4)3.
9–11

4

Owing to the large three-dimensional interstitial spaces in their structures, anti-
NASICON Fe2(MO4)3 (M = Mo or W) are promising intercalation electrode materials. The
Earth abundant, non-toxic, and inexpensive characteristics of iron makes Fe2(MoO4)3 a
potential candidate for rechargeable sodium or lithium ion batteries.
12–14
Fe2(MoO4)3 has
an open 3-D network structure adopting the monoclinic, P21/a space group, as well as a
high temperature polymorph with an orthorhombic crystal structure. Structural
transformation from monoclinic to orthorhombic polymorph takes place at 518 ˚C. The
tungstate analogue, Fe2(WoO4)3, adopts the same monoclinic structure, however, it
disproportionates into Fe2WO6 and WoO4 at 700 ˚C.
15,16

In 1985, Reiff and Zhang first reported topochemical lithium insertion into
Fe2(MoO4)3.
17
Chemical lithium insertion into Fe2(MoO4)3 was performed by reacting
ferric molybdate with LiI in dry acetonitrile for 3 weeks. In 1986, Torardi and Prince
reported the structure of chemically lithium-inserted iron molybdate [i.e., Li2Fe2(MoO4)3]
by neutron diffraction experiments.
18
Li2Fe2(MoO4)3 was shown to crystallize into an
orthorhombic structure with Pbcn space group and lithium ions occupy tetrahedral
interstitial sites in the framework. The interstitial sites occupied by the lithium ions and the
conduction channels are both generated along the same c-axis. Thus, the lithium ions can
easily move through these conduction channels.
Lithium can be inserted into the Fe2(MO4)3 (M = Mo or W) frameworks both
electrochemically and chemically. In 1987, Manthiram and Goodenough demonstrated
electrochemical lithium insertion into both iron molybdate and iron tungstate.
11
Although
up to eight lithium ions can be inserted per formula unit of Fe2(MoO4)3, the insertion is
topotactic only when less than 2.0 lithium ions per formula unit is inserted. Electrochemical

5

insertion results in two-phase structural change as lithium ions are inserted over the range
of 0 to 1.7 lithium ions per formula unit and the monoclinic pristine structure transforms
into orthorhombic lithium-inserted one. These lithium-inserted orthorhombic (Pbcn)
structures further transform into another orthorhombic structure (Pnma) at 537 K and 579
K for molybdenum and tungsten analogues, respectively. Moreover, the Fe2(MO4)3
frameworks can undergo topotactic insertion/extraction of lithium ions owing to large
vacant spaces in the three-dimensional framework. The use of large (MO4)3
2-
anions
permits fast lithium ion insertion/extraction in this open three-dimensional framework and
stabilizes the redox potentials Fe
3+
/Fe
2+
that give open-circuit voltages Voc > 2.5 V as well
as allow access to Fe
2+
/Fe
+
couple.
11

Although there is an agreement on the two-phase mechanism for lithium insertion into
Fe2(MO4)3, previous reports on sodium insertion into the same framework are
contradictory. In 1984, Nadiri et al. first investigated Na
+
-insertion into monoclinic
Fe2(MO4)3 and reported single-phase solid-solution type mechanism over the ranges 0.3 ≤
x ≤ 1.0 and 1.10 ≤ x ≤ 1.60 (i.e., x is inserted sodium ion per formula unit).
19
In 1990,
sodium insertion into anti-NASICON Fe2(MO4)3 (M = Mo or W) frameworks has also been
investigated by Bruce and Miln.
20
Electrochemical cycling data indicated a plateau-type
constant cell voltage upon sodium insertion into these frameworks suggesting a two-phase
structural change in the range of 0 ≤ x ≤ 2.0. Recently, Heo et al. reported the structure of
sodiated iron molybdate (i.e., Na2Fe2(MO4)3) by ab initio structure determination from
powder diffraction data.
21
Four distinct sodium positions in the monoclinic unit cell with
P21/a space group were found by difference Fourier synthesis maps (Figure 1.2).

6

Figure 1.2. The (010) view of the structure of fully sodiated Na2Fe2(MoO4)3. Reprinted
with permission from ref. 21 (Copyright 2018 American Chemical Society).

These studies demonstrate the investigation of lithium and sodium insertion into
Fe2(MO4)3 (M = Mo or W) frameworks to some extent, however, the key factors driving
the two-phase and single-phase structural change behaviors upon Li
+
and Na
+
insertion has
not been very well understood yet. Recently, Yue et al. have investigated the mechanism
of lithium and sodium intercalation into orthorhombic polymorph of Fe2(MoO4)3 by
synchrotron-based ex situ X-ray diffraction, X-ray absorption spectroscopy and electron
microscopy, along with density functional theory calculations.
22
A two-phase structural
change upon lithium insertion was demonstrated, which occurs with a discrete occupation
path. Whereas, sodium insertion is a single-phase mechanism governing by the pseudo-
continuous filling of sodium sites in the framework (Figure 1.3).

7

Figure 1.3. Comparison between sodiation and lithiation process in Fe2(MoO4)3.
Schematic diagrams of ‘‘discrete occupation’’ and ‘‘pseudocontinuous occupation’’ during
Li and Na ions intercalation into Fe2(MoO4)3. Solid red circles and dash green ellipses
stand for Li1 (or Na1) and Fe2(MoO4)3 frameworks, respectively. Reprinted with
permission from ref. 22 (Copyright 2015 Springer Nature Limited).

In chapter 2 and 3, the mechanism of reversible lithium and sodium intercalation into
anti-NASICON type monoclinic Fe2(MO4)3 (M = Mo or W) frameworks will be presented
with an emphasis on the reorganization of the rigid polyhedral units within the structure
framework. The correlation of structural changes in the frameworks of these polyanionic
cathode materials with their electrochemical properties will be demonstrated through a
wide spectrum of structural and electrochemical characterization techniques.

8

1.3. Binary nickel sulfide system for energy applications

Among the family of metal sulfides, nickel sulfide is an appealing class of energy
materials that possess interesting structural, compositional, and magnetic properties.
Owing to the relatively complex nickel-sulfur phase diagram, binary nickel sulfide system
consists of several compositions (e.g., Ni3S2, Ni7S6, Ni9S8, NiS, Ni3S4, NiS2) (Figure 1.4).
23

According to the phase diagram, heazlewoodite Ni3S2 is the most metal-rich phase of the
binary Ni-S system that displays Pauli paramagnetic behavior.
24
The low-temperature
polymorph of Ni3S2 possess a rhombohedral structure that is stable up to 843 K. Above 843
K, it transforms into a high-temperature, non-stoichiometric Ni3-xS2 polymorph.
24
Due to
the continuous network of Ni-Ni bonds thought the structure, heazlewoodite Ni3S2
demonstrates intrinsic metallic behavior.
25
Earth abundant elements, low-cost, and high
electrical conductivity make Ni3S2 is a promising candidate for various electrochemical
applications. Ni3S2 has been shown to be an electrocatalytically active material for both
hydrogen evolution reaction (HER) and oxygen evolution reaction (OER) of water
splitting.
26–29
Besides electrocatalytic water splitting, heazlewoodite Ni3S2 has also been
used as electrodes and electroactive materials for lithium-sulfur batteries
30
and
supercapacitors
31–33
, respectively.
Godlevskite Ni9S8 is another metal-excess non-stoichiometric form of nickel sulfide
with an orthorhombic structure. In 2011, Tatsumisago and coworkers reported a colloidal
synthesis of Ni9S8 nanorods by thermal decomposition of nickel acetylacetonate in a mixed
solution of 1-dodecanethiol and oleylamine at 280 ˚C.
34
More recently, colloidal synthesis
of two dimensional cross-like Ni9S8 nanoparticles were also demonstrated. In the synthesis,
halide ions were employed a reactivity-directing agents to slow down the growth kinetics

9

of the nanocrystals.
35
Similar to the most metal-rich Ni3S2, Ni9S8 has also been shown to
be an active material for the hydrogen evolution reaction
36
, oxygen evolution reaction,
37

and supercapacitors,
38
as well as an electrode material for all-solid-state lithium secondary
batteries
34
.

Figure 1.4. Calculated condensed nickel-sulfur phase diagram by Waldner and Pelton.
Reprinted with permission from ref. 23 (Copyright 2004 Deutsche Gesellschaft für
Materialkunde).

The stoichiometric form of nickel sulfide has two polymorphs: the low-temperature
rhombohedral β-NiS (millerite) and high-temperature hexagonal α-NiS. The β-to-α phase
transition occurs at 379 ˚C resulting in a 4% volume change.
39
High-temperature α-NiS has
been shown to exhibit metal-to-semiconductor and paramagnetic-antiferromagnetic
transitions.
40
Sulfur-rich, non-stoichiometric Ni3S4 (polydymite) is a low-temperature
phase with a cubic spinel crystal structure that disproportionates into NiS and NiS1.03 at
high temperatures.
41
Conventionally, Ni3S4 has been prepared along with NiS and NiS2

10

impurities by a solid-state reaction between elemental nickel and sulfur at 300 ˚C for 8
months because its instability at higher temperatures (above 360 ˚C).
42
Ni3S4 exhibits
higher electrolyte resistance than those of other binary nickel sulfide phases when it is
placed on n-doped graphene.
43
Similar to metal-rich nickel sulfides (i.e., Ni3S2 and Ni9S8),
stochiometric and sulfur-rich nickel sulfides (α-NiS, β-NiS and Ni3S4) have been
commonly used as electrode materials in a number of energy applications including
supercapacitors
44–47
, rechargeable ion batteries,
48–51
and electrocatalysis
52,53
.
As Earth abundant, low-cost, rich-valence state, and environmentally friendly
materials, various compositions of binary nickel sulfides have been shown to be promising
electrocatalysts for hydrogen evolution and oxygen evolutions reactions. However, for the
direct comparison of the impact of different compositions and phases of binary nickel
sulfides on their intrinsic catalytic activities, these materials must be synthesized and
analyzed under identical conditions. For this purpose, Jiang et al. recently reported
preparation and electrochemical characterization of three different compositions of nickel
sulfides under the same conditions.
54
Ni3S2, α-NiS and NiS2 were synthesized though a
microwave-assisted solvothermal approach by the reaction of nickel acetate and elemental
sulfur in oleylamine. By adjusting the amount of elemental sulfur in the reaction, highly
crystalline trigonal Ni3S2, hexagonal α-NiS, and cubic NiS2 were obtained (Figure 1.5).
Electrocatalytic performance of these phases toward hydrogen evolution reaction was
examined under alkaline conditions. Linear sweep voltammetry measurements indicate
that Ni3S2 has the lowest overpotential of -355 mV to reach a current density of 10 mA cm
-
2
, which is followed by NiS2 and α-NiS with overpotentials of -454 mV and -474 mV,
respectively. These differences in the intrinsic catalytic activities of nickel sulfides were

11

attributed to higher electrical conductivity and large electrochemically active surface are
of Ni3S2 based on theoretical calculations and electrochemical impedance spectroscopy
measurements.

Figure 1.5. Figure 1.5. Unit cell structures, scanning electron microscopy images, and
XRD patterns of (a) NiS, (b) NiS2, and (c) Ni3S2. Reprinted with permission from ref. 54
(Copyright 2016 Royal Society of Chemistry).

1.4. Solution-phase synthesis of binary nickel sulfides

Conventional methods to prepare binary nickel sulfides include solid-state reaction of
elemental nickel and sulfur at high temperatures in an evacuated silica tube.
55,56
Solid-state
or vapor-phase reactions are usually carried out under reducing conditions in a stream of
H2/H2S.
57,58
Several other strategies including hydrolysis-controlled precipitation
59
,
hydrothermal
60
and solvothermal approaches
61
, in situ growth on a nickel substrate
62
, ionic

12

liquid assisted sonochemical method
63
and one-step electrodeposition approach
64
have
been widely adopted for the synthesis of micro- and nanostructures of nickel sulfides.
These synthetic approaches generally yield products with large particle size and/or ill-
defined morphologies. However, for specific applications in electrocatalysis or
rechargeable ion batteries, precise and simultaneous control of composition, phase, and
morphology is crucial. For instance, the nanoscale synthesis of nickel sulfides would be
beneficial for rechargeable ion battery applications by decreasing the diffusion length of
ions. Similarly, increasing surface area of nanoparticles can potentially lead to an increase
in their catalytic activity for electrocatalytic water splitting applications. In this regard,
solution-phase synthesis of colloidal nickel sulfide nanoparticles with well-defined
morphologies has been extensively explored.
The synthetic methods previously reported for the preparation of nanoscale nickel
sulfides include solventless thermolytic decomposition,
65
solution-phase arrested
precipitation,
66
thermal decomposition of nickel acetylacetonate in the presence of 1-
dodecanethiol,
67
and decomposition of nickel dithiocarbamates,
68–70
alkyl xanthates,
71

thiobiurets,
72
2-mercaptobenzothiazole complexes,
73
(TMEDA)Ni(SCOC6H5)2,
74
and
nickel polysulfide complexes
75
as single source precursors in high-boiling solvents. In
2004, Korgel et al. demonstrated high degree of morphology control over β-NiS
nanocrystals through solventless thermolytic decomposition of single source nickel thiolate
precursors in octanoate.
65
Although the resulting nanorods and triangular nanoprisms
display narrow size and shape distributions, cubic Ni3S4 was also observed as an impurity
phase in the form of misshapen needles and particulates. More recently, Batteas and
coworkers reported shape-controlled synthesis of cubic Ni3S4 by a one-pot colloidal

13

synthesis.
76
In a typical synthesis, nickel salts (i.e., NiCl2∙H2O or Ni(acac)2(H2O), where
acac = CH3COCH2COCH3) were dissolved in a mixture of high-boiling solvents and
ligands (i.e., 1-octadecene, oleylamine, oleic acid and 1,2 hexadecanediol), and 1-
dodecanthiol was injected as sulfur source. Nanopyramids and nanoprisms of Ni3S4 were
obtained with high crystallinity and narrow size distributions (Figure 1.6). Nickel
precursors and capping ligands have been shown to be the main factors that dictate
morphology of the nanoparticles.

Figure 1.6. (a) TEM image of the Ni3S4 nanoprism synthesized using NiCl2 as the Ni
precursor; (b) TEM image of the Ni3S4 tetrahedron (nanopyramids) synthesized using
Ni(acac)2 as the Ni precursor; (c and d) length (c) and width (d) distribution for the Ni 3S4
nanoprisms (40 nm × 10 nm); (e) length distribution of the edges for the Ni3S4 tetrahedra
(16 nm). Reprinted with permission from ref. 76 (Copyright 2014 Royal Society of
Chemistry).

14

Distinct phases of binary nickel sulfide system have also been synthesized with various
morphologies including micrometer-size hollow spheres and porous sponge-like
nanostructures of heazlewoodite Ni3S2 by a γ-irradiation
77
and biomolecule assisted
routes
78
, respectively; nanowhiskers of millerite β-NiS on nickel foils as well as its 3D
nanostructures by a general solution-phase
79
and solvothermal
80
approaches, respectively;
layered rolled structures of hexagonal α-NiS through a low-temperature solution-phase
route
81
; 1D rod-like and 2D cross-like nanostructures of godlevskite Ni9S8 via thermal
decomposition of nickel acetylacetonate
34
and heating up method in the presence of halide
ions
35
, respectively.
Synthetic control in that binary nickel-sulfur phase space has also been explored to
some degree. In 2001, Jeong and Manthiram investigated the phase-controlled synthesis of
nickel sulfides through the reaction of nickel chloride with sodium dithionate in aqueous
solutions under ambient conditions.
82
It was found that compositions and structures of
nickel sulfides can be tuned by adjusting pH of the reaction and amounts of the precursors.
Highly basic (pH≥7) and highly acidic (pH≤3) conditions yield poorly crystalline nickel
sulfides or crystalline sulfur, however, at moderate conditions (3 ≤ pH ≤ 6) crystalline
nickel sulfides with various compositions can be obtained (Figure 1.7).

15

Figure 1.7 Diagrammatic representation of the compositions obtained with 50 mL of 0.25
M NiCl2•6H2O (12.5 mmol) and 25-300 mL of 1 M sodium dithionite (25-300 mmol) at
various pH values. The compositions within the rectangular boxes are crystalline Ni3Sx
while those in the shaded area are crystalline sulfur (dark area) and amorphous or poorly
crystalline NiySx (light area). Reprinted with permission from ref. 82 (Copyright 2001
American Chemical Society)

Recently, decomposition of single-source precursors in high boiling solvents have been
widely investigated for the phase-controlled synthesis of binary nickel sulfides in the
nanoscale regime. In 2016, Leeuw et al. prepared square-planar nickel bis(thiocarbamate)
complexes [i.e., Ni(S2CNR2)2] and employed as single source reactants for the synthesis of
nickel sulfide nanoparticulates.
68
At low temperatures (150 ˚C), decomposition of isobutyl
derived nickel bis(thiocarbamate) leads to formation of the hexagonal α-NiS phase.
Whereas, at higher temperatures (280 ˚C), phase-pure rhombohedral β-NiS
nanoparticulates were obtained under otherwise identical conditions. When thiuram
disulfide [i.e., (
i
Bu2NCS2)2] was added to the decomposition mixture, formation of the

16

cubic Ni3S4 phase along with hexagonal α-NiS was observed at low temperatures (150 –
180 ˚C). In addition to reaction temperature, precursor concentration was identified as an
important synthetic lever to control phase of the nanocrystals and with an increasing
precursor concentration, increasing proportions of the cubic Ni3S4 phase was noted.
In a very similar study, Gervas et al. demonstrated synthetic phase control over nickel
sulfide nanostructures by using two different nickel complexes as single source
precursors.
70
Piperidine and tetrahydroquinoline dithiocarbamate ligands, and the
corresponding nickel complexes, were prepared and nickel sulfide nanoparticulates were
synthesized by the solvothermal decomposition of these complexes in several primary
amine coordinating solvents. Reaction temperature and primary amine solvents were found
to be phase-directing factors. At 230 ˚C, cubic Ni3S4 and rhombohedral Ni3S2 phases were
obtained in the presence of dodecylamine and hexadecylamine, respectively. Although
these studies establish the influences of various reaction parameters on the control of nickel
sulfide phases, neither of them has sufficient control over the morphology, and the as-
synthesized nanocrystals were polydispersed with ill-defined shapes. In chapter 4, a high
degree of control over nanocrystal phase in the complicated binary nickel-sulfur phase
space, and preparation of morphologically well-defined binary nickel sulfide nanocrystals,
will be demonstrated and discussed.

17

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24

Chapter 2. Influence of Rotational Distortions on Li
+
- and Na
+
-Intercalation in Anti-
NASICON Fe2(MoO4)3*

*Published in Chem. Mater. 2016, 28, 4492-4500.

2.1. Abstract

Anti-NASICON Fe2(MoO4)3 (P21/c) shows significant structural and electrochemical
differences in the intercalation of Li
+
and Na
+
ions. To understand the origin of this
behavior, we have used a combination of in situ X-ray and high-resolution neutron
diffraction, total scattering, electrochemical measurements, density functional theory
calculations, and symmetry-mode analysis. We find that for Li
+
-intercalation, which
proceeds via a two-phase monoclinic-to-orthorhombic (Pbcn) phase transition, the host
lattice undergoes a concerted rotation of rigid polyhedral subunits driven by strong
interactions with the Li
+
ions, leading to an ordered lithium arrangement. Na
+
-intercalation,
which proceeds via two solid solution regions of the monoclinic structure, similarly
produces rotations of the lattice polyhedral subunits. However, using a combination of total
neutron scattering data and density-functional theory calculations, we find that while these
rotational distortions upon Na
+
intercalation are fundamentally the same as for Li
+
intercalation, they result in a far less coherent final structure, with this difference attributed
to the substantial difference between the ionic radii of the two alkali metals.

2.2. Introduction

The demand for rechargeable lithium- or sodium-ion batteries will continue to grow as
renewable energy sources are integrated into the electrical grid, and electric vehicles
continue to be developed.
1–3
In rechargeable ion batteries, positive electrode materials

25

undergo reversible intercalation, displacement, or conversion processes to store
electricity.
4,5
Present state-of-the-art positive intercalation electrodes for Li
+
or Na
+
ion
batteries rely on either layered oxides, spinel oxides, or polyanionic compounds.
6

Polyanionic compounds have structural frameworks composed of corner-sharing
polyhedral subunits. This allows much greater lattice flexibility than in layered or spinel
oxides and contributes to noteworthy structural stability with respect to electrochemical
cycling. Polyanionic compounds are also considered to be safer than alternative oxides,
because they do not release oxygen upon decomposition, which can exacerbate thermal
runaway during cell failure.
7,8
Because the utility of polyanionic cathodes is linked to the
intrinsic flexibility of their host lattice, the development of next-generation polyanionic
electrodes requires understanding the relevant structural distortions that occur across
multiple polyanionic compounds as Li
+
and Na
+
are inserted.
NASICON-related structures, with general formula MM’(XO4)3, were among the first
polyanionic compounds investigated as intercalation electrodes.
7
The structures feature
large interstitial voids formed by three-dimensionally connected polyhedra, creating a
robust structural topology. The chemical versatility of these materials structures means
substitution of the polyanionic group can be exploited to tune the position of the redox
couples.
9,10
Anti-NASICON Fe2(MoO4)3 is a promising cathode for both Li
+
and Na
+
ion
cathodes, owing to its highly reversible accommodation of intercalated ions.
11–14

Fe2(MoO4)3 crystallizes in a thermodynamically preferred, low-temperature monoclinic
(P21/c) structure that consists of FeO6 octahedra and MoO4 tetrahedra interconnected
through corner sharing oxygen atoms (Figure 2.1). The basic motif of the structure is
described as the “lantern unit”, composed of three MoO4 tetrahedra that connect two FeO6

26

octahedra (Figure 2.1b). In the anti-NASICON structure, these units stack in an antiparallel
fashion along the (2b + c). In contrast, in the NASICON structure equivalent units stack in
a parallel fashion along the c-axis.
7

Figure 2.1. Illustration of (a) the unit cell of pristine, monoclinic Fe2(MoO4)3 and (b) a
“lantern unit” that consists of three MoO4 tetrahedra connecting two FeO6 octahedra.

While the electrochemical performance of Li
+
and Na
+
(de)insertion into Fe2(MoO4)3
has been studied previously, the existing knowledge in the literature is often contradictory,
and the mechanism of structural transformations from the parent phase to the intercalated
phase remains unknown for monoclinic Fe2(MoO4)3. Nadiri et al. reported a single-phase
structural change in monoclinic Fe2(MoO4)3 during Na
+
(de)insertion in the composition
ranges of 0.3 ≤ x ≤ 1.0 and 1.10 ≤ x ≤ 1.60 for NaxFe2(MoO4)3.
12
In contrast, Bruce et al.
reported the presence of two phases throughout the composition range of 0 < x < 2 through
ex-situ X-ray diffraction of powder samples that were electrochemically intercalated.
13

27

Recent studies by Yue et al. reported a two-phase structural change for Li
+
(de)insertion,
but a single-phase process for Na
+
-intercalation in the high-temperature, orthorhombic (not
monoclinic) polymorph of Fe2(MoO4)3.
15
To reconcile these various conflicting reports, it
is crucial to understand the mechanism by which the structural framework transforms in
order to accommodate the intercalation of guest ions.
Herein, we present an investigation into the different insertion behaviors of Li
+
and Na
+

in monoclinic Fe2(MoO4)3, with an emphasis on the reorganization of the rigid polyhedral
units within the structure framework upon insertion. We have used a combination of
structural and electrochemical tools, coupled with an analysis of symmetry-allowed
distortions, to characterize the relationship between the structural evolution of Fe2(MoO4)3
during (de)insertion of Na
+
and Li
+
and their corresponding electrochemical

properties. We
find the insertion processes for Li
+
and Na
+
differ in the ability of rigid MoO4 tetrahedra to
rotate in order to accommodate the presence of guest ions within the lattice framework.
This mechanism, whereby intercalation proceeds through cooperative rotations, is found
to be strongly influenced by the size and nature of the alkali ion. This result is therefore
also relevant to other polyanionic electrodes and, more broadly, to intercalation into any
system containing a mixture of rigid and soft bonding units.

2.3. Results and Discussion

Pristine Fe2(MoO4)3 was prepared by a precipitation method following the approach of
Peng et al.
16
After annealing at 400
°
C for 6 h, the phase pure powder possessed a nanorod-
like morphology (Figure 2.2). Figure 2.3 shows the results of the Rietveld refinement of
the monoclinic P21/c structure against the synchrotron X-ray (Figure 2.3a) and neutron

28

diffraction (Figure 2.3b) patterns of the as-prepared Fe2(MoO4)3 electrode material. The
diffraction patterns were refined simultaneously with structural parameters starting from
the model of Chen et al.
17
Detailed structural parameters including atomic positions, cell
parameters, and anisotropic displacement parameters are given Table 2.1. Both diffraction
patterns could be fully indexed to lattice planes from the expected monoclinic phase
without any evidence of impurity peaks, confirming that the parent Fe2(MoO4)3 utilized
throughout this study was phase pure.

Figure 2.2. TEM image of pristine Fe2(MoO4)3.

29

Figure 2.3. Rietveld refinements of pristine, monoclinic Fe2(MoO4)3 against (a)
synchrotron XRD pattern obtained at the 11-BM beamline at Argonne National Laboratory
and (b) neutron diffraction pattern obtained at the POWGEN beamline at Oak Ridge
National Laboratory. The weighted profile R-factor (R wp) was determined to be 6.01% for
all banks being refined simultaneously and weighed equally. Absorption correction was
carried out using absorption function 0 in GSAS. Only the diffraction pattern from bank 3
at POWGEN is given here.

Table 2.1. Resulting unit cell, atomic positions and atomic displacement parameters from
the Rietveld refinement of Fe2(MoO4)3 against synchrotron X-ray powder diffraction data
and neutron diffraction data (the Rwp was determined to be 6.01% for all banks being refined
simultaneously). Both sets of data were collected at room temperature.

Space group P2 1/c
a (˚A) 15.73802(26)
b (˚A) 9.23523(13)
c (˚A) 15.70359(25)
β (deg) 109.1209(31)

30

Atom Wyckoff Occupancy x y z Uiso (Å
2
)
Mo1 4e 1.0 -0.4874(7) 0.24655(13) -0.49453(5) 0.0443(5)
Mo2 4e 1.0 -0.13732(7) 0.12818(10) 0.22359(7) 0.0443(5)
Mo3 4e 1.0 -0.24898(7) 0.11241(10) -0.10977(5) 0.0443(5)
Mo4 4e 1.0 -0.38521(7) 0.61604(10) -0.23417(7) 0.0443(5)
Mo5 4e 1.0 -0.21211(5) 0.63085(10) 0.13911(5) 0.0443(5)
Mo6 4e 1.0 -0.01982(7) 0.74458(11) -0.01712(5) 0.0443(5)
Fe1 4e 1.0 -0.31899(5) 0.9651(9) 0.06167(5) 0.06116(31)
Fe2 4e 1.0 -0.04482(5) 0.4593(8) 0.32522(5) 0.06116(31)
Fe3 4e 1.0 -0.18524(5) 0.4738(10) -0.06614(5) 0.06116(31)
Fe4 4e 1.0 -0.41592(5) 0.983(8) -0.31004(5) 0.06116(31)
O1 4e 1.0 0.0001(7) 0.3872(13) 0.5818(7) 0.03967(22)
O2 4e 1.0 -0.1715(8) 0.4075(13) 0.8167(8) 0.03967(22)
O3 4e 1.0 -0.0949(7) 0.1937(11) 0.7331(7) 0.03967(22)
O4 4e 1.0 -0.0544(7) 0.4974(13) 0.72(7) 0.03967(22)
O5 4e 1.0 -0.1443(7) 0.4253(11) 0.3782(7) 0.03967(22)
O6 4e 1.0 -0.2728(7) 0.5194(13) 0.4593(7) 0.03967(22)
O7 4e 1.0 -0.4139(8) 0.1099(13) 0.0065(8) 0.03967(22)
O8 4e 1.0 -0.2464(7) 0.2854(10) -0.0621(7) 0.03967(22)
O9 4e 1.0 -0.4437(7) 0.3536(11) 0.1045(7) 0.03967(22)
O10 4e 1.0 -0.9753(8) 0.3088(13) -0.5876(7) 0.03967(22)
O11 4e 1.0 -0.0674(7) 0.3818(10) -0.003(7) 0.03967(22)
O12 4e 1.0 -0.501(7) 0.3627(11) -0.0895(7) 0.03967(22)
O13 4e 1.0 -0.2325(8) 0.3829(13) 0.625(7) 0.03967(22)
O14 4e 1.0 -0.5128(8) 0.0396(11) -0.2644(7) 0.03967(22)
O15 4e 1.0 -0.3324(7) 0.1065(13) -0.216(7) 0.03967(22)
O16 4e 1.0 -0.3626(7) 0.9432(13) 0.1701(8) 0.03967(22)
O17 4e 1.0 -0.1926(8) 0.9597(13) 0.5496(7) 0.03967(22)
O18 4e 1.0 -0.3007(8) 0.9378(10) 0.3647(8) 0.03967(22)
O19 4e 1.0 -0.3116(7) 0.9339(13) 0.6503(7) 0.03967(22)
O20 4e 1.0 -0.5939(8) 0.3111(13) -0.4994(8) 0.03967(22)
O21 4e 1.0 -0.3913(7) 0.8062(13) -0.2345(8) 0.03967(22)
O22 4e 1.0 -0.121(8) 0.6577(13) -0.0759(7) 0.03967(22)
O23 4e 1.0 -0.1228(7) 0.6027(11) 0.2401(8) 0.03967(22)
O24 4e 1.0 -0.2257(8) 0.8171(11) 0.1181(7) 0.03967(22)

31

Figure 2.4. Raman spectra of (a) pristine Fe2(MoO4)3 under ambient conditions and (b)
Li
+
-inserted and Na
+
-inserted Fe2(MoO4)3 along with the pristine Fe2(MoO4)3 in the air-
free quartz cell. *Indicates band corresponding to air-free cell.

Raman spectroscopy showed that the as-prepared Fe2(MoO4)3 possessed bands
characteristic of the monoclinic structure at 988, 967, 930 cm
–1
(symmetric stretching
modes of terminal Mo=O bonds in three distinct MoO4 tetrahedra); 817, 776 cm
–1

(asymmetric stretching modes of MoO4 units); and 356 cm
–1
(MoO4 bending mode) (Figure
2.4).
18
X-ray photoelectron spectra (XPS) of the Fe2p and Mo3d regions suggest there is
only one chemical environment for both Fe and Mo in the pristine, as-prepared Fe2(MoO4)3
by the presence of characteristic Fe
3+
2p3/2 and Mo
6+
3d5/2 peaks with binding energies of
711.3 eV and 232.5 eV, respectively (Figure 2.5).
19,20

32

Figure 2.5. High-resolution XPS spectra of pristine Fe2(MoO4)3.

2.3.1. Li
+
insertion
The first ten galvanostatic electrochemical cycles of Fe2(MoO4)3 against Li/Li
+
at a
current rate of C/10 are given in Figure 2.6a. The initial specific capacity of 92 mA·h·g
−1

agrees well with the theoretical capacity of intercalation of two Li
+
ions per formula unit
(90 mA·h·g
−1
), with the slight excess capacity likely associated with the formation of a
passivating solid electrolyte interface layer.
21
Without rigorous optimization of the cell
assembly, the batteries showed a high capacity retention of around 90% of the initial
capacity after 25 cycles. The derivative of the galvanostatic cycling (inset of Figure 2.6a)
shows two peaks centered at 3.0 V and 3.05 V during reduction and oxidation, respectively.
Such a small polarization is reflective of the highly reversible intercalation of Li
+
into the
framework of Fe2(MoO4)3. The voltage-composition curve exhibits a single plateau over a
wide range of lithium content during Li
+
(de)insertion, which indicates the intercalation
process predominantly occurs through a two-phase process where a lithium-rich phase is
created directly rather than a solid-solution process, since there is a continuous change in
the lithium content throughout each particle.
11

33

Figure 2.6. (a) Galvanostatic electrochemical cycling of Fe2(MoO4)3 against Li
+
insertion
and its derivative (shown as inset). (b) 2D pattern based on the in situ XRD of Li
+
insertion
into Fe2(MoO4)3.

The coexistence of two phases was further demonstrated using the in situ X-ray
diffraction, taken continuously during electrochemical cycling, and is shown as a heat map
in Figure 2.6b with the y-axis corresponding to the equivalents of inserted Li
+
. During the
discharging process, new reflections at 24.5°, along with several others at 36° and 38.5°,
clearly appear (Figures 2.7 and 2.8). The intensity of other reflections (e.g., at 23.5°, 36.8°,
and 39.8°) gradually decreases until finally disappearing as the Li
+
content increases. As
shown in Figure 2.8, all reflections of the fully electrochemically lithiated phase, except

34

those contributed by the beryllium cell used for the in situ measurements, can be indexed
to the orthorhombic Li2Fe2(MoO4)3 phase originally reported by Torardi et al.
22

During the charging process, changes in the intensities of reflections were observed
exactly in the opposite way, indicating excellent structural reversibility upon Li
+

(de)insertion. Positions of the reflections do not vary with Li
+
insertion in the pristine
Fe2(MoO4)3, which excludes the possibility of a solid solution process. These changes in
the intensities of reflections are illustrative of the evolution of the initial monoclinic
Fe2(MoO4)3 structure into orthorhombic Li2Fe2(MoO4)3 with the increasing Li
+
content,
and the coexistence of pristine and lithiated structures in various ratios depending on the
Li
+
concentration. Thus, a two-phase Li
+
(de)insertion process in the anti-NASICON
Fe2(MoO4)3 was confirmed for the first time through in situ XRD.

Figure 2.7. Rietveld refinement of one-hour X-ray diffraction pattern on pristine
Fe2(MoO4)3.

35

Figure 2.8. Rietveld refinement of one-hour X-ray diffraction pattern on fully
electrochemically lithiated Li2Fe2(MoO4)3. Intensities offset at 30, 45, 47, and 49 degrees
are contributed by the beryllium cell used for in situ X-ray diffraction.

The crystal structure of orthorhombic Li2Fe2(MoO4)3 (Figure 2.9) shows an identical
structural topology to the parent monoclinic phase with regards to polyhedral connectivity,
so the transformation is purely displacive in nature. While the in situ XRD experiments
provide some insight into the mechanism for these displacements, the patterns do not have
sufficient resolution or intensity to precisely refine changes in the atomic positions.
Therefore, Fe2(MoO4)3 was chemically lithiated as discussed in the experimental section.
The results of the Rietveld refinement of the structure against synchrotron X-ray and
neutron diffraction data is shown in Figure 2.9 with detailed structural parameters listed in
Table 2.2.

36

S2.
Figure 2.9. Crystal structure of (a) pristine, monoclinic Fe2(MoO4)3 and (b) fully lithiated,
orthorhombic Li2Fe2(MoO4)3.

Figure 2.10. Refinement results of neutron powder diffraction patterns of chemically
lithiated Li2Fe2(MoO4)3.

37

Table 2.2. Resulting unit cell, atomic positions, and atomic displacement parameters from
the Rietveld refinement of lithiated Li2Fe2(MoO4)3 against neutron diffraction data (the Rwp
was determined to be 3.02% for all banks being refined simultaneously).

Space group Pbcn
a (˚A) 12.89461(7)
b (˚A) 9.48373(6)
c (˚A) 9.33841(6)
Atom Wyckoff Occupancy x y z Uiso (˚A2)
Fe1 8d 1.0 0.3795(14) 0.2453(17) 0.4672(17) 0.04275 (32)
Mo1 4c 1.0 0 0.463(28) 0.25 0.038 (5)
Mo2 8d 1.0 0.3573(18) 0.4027(20) 0.1046(20) 0.038(5)
O1 8d 1.0 0.1627(14) 0.0832(27) 0.1087(26) 0.02533 (25)
O2 8d 1.0 0.0941(21) 0.3531(24) 0.165618) 0.02533 (25)
O3 8d 1.0 0.2773(16) 0.3299(20) 0.9712(4) 0.02533 (25)
O4 8d 1.0 0.435(11) 0.0698(26) 0.373(18) 0.02533 (25)
O5 8d 1.0 0.487(21) 0.356(21) 0.0698(18) 0.02533 (25)
O6 8d 1.0 0.3218(11) 0.3261(19) 0.2733(30) 0.02533 (25)
Li1 8d 1.0 0.1875(5) 0.2239(7) 0.277(13) 0.01267 (10)

To describe the transition between the monoclinic and orthorhombic structures in terms
of the reversible symmetry-allowed distortions, the AMPLIMODES program at the Bilbao
Crystallographic Server
23
was used to identify the transformation matrix (Table 2.3 and
2.4) for the transition from the low-symmetry (P21/c) pristine structure to the higher-
symmetry (Pbcn) lithiated structure. The vectors corresponding to the direction of the
displacements for each element were calculated and are illustrated in Figure 2.11. A
complete list of the vectors is given in Tables 2.5 and 2.6.

38

Table 2.3. Transformation matrix for converting parent structure (Pbcn) into reference
structure (P21/c) for symmetry mode analysis.

[ 1 0 -1 ] [ 1 ]
[ 1 0 1 ] [ 1/2 ]
[ 0 -1 0 ] [ 0 ]

Table 2.4. Unit cell parameters for parent structure in reference cell settings.

Space Group P21/c
a (Å) 16.006628
b (Å) 9.338410
c (Å) 16.006628
β (
o
) 107.332512

39

Table 2.5. Atomic pairing for the parent phase and the distorted phase in the reference
settings.

Atom Wyckoff x y z Atom Wyckoff x y z
Fe 1 4e 0.5624 0.5328 0.1829 Fe 4 4e -0.41592 0.983 -0.31004
Fe 1_2 4e 0.1829 0.9672 0.5624 Fe 3 4e -0.18524 0.4738 -0.06614
Fe 1_3 4e 0.9376 0.4672 0.3171 Fe 2 4e -0.04482 0.4593 0.32522
Fe 1_4 4e 0.3171 0.0328 0.9376 Fe 1 4e -0.31899 0.9651 0.06167
Mo 1 4e 0.4815 0.75 0.4815 Mo 1 4e -0.4874 0.24655 -0.49453
Mo 1_2 4e 0.0185 0.25 0.0185 Mo 6 4e -0.01982 0.74458 -0.01712
Mo 2 4e 0.63 0.8954 0.2727 Mo 4 4e -0.38521 0.61604 -0.23417
Mo 2_2 4e 0.2727 0.6046 0.63 Mo 3 4e -0.24898 0.11241 -0.10977
Mo 2_3 4e 0.87 0.1046 0.2273 Mo 2 4e -0.13732 0.12818 0.22359
Mo 2_4 4e 0.2273 0.3954 0.87 Mo 5 4e -0.21211 0.63085 0.13911
O 1 4e 0.37295 0.8913 0.21025 O 15 4e 0.3324 0.8935 0.216
O 1_2 4e 0.21025 0.6087 0.37295 O 13 4e 0.2325 0.6171 0.375
O 1_3 4e 0.12705 0.1087 0.28975 O 23 4e 0.1228 0.1027 0.2599
O 1_4 4e 0.28975 0.3913 0.12705 O 18 4e 0.3007 0.4378 0.1353
O 2 4e 0.4736 0.8344 0.3795 O 9 4e 0.4437 0.8536 0.3955
O 2_2 4e 0.3795 0.6656 0.4736 O 7 4e 0.4139 0.6099 0.4935
O 2_3 4e 0.0264 0.1656 0.1205 O 1 4e 0 0.1128 0.0818
O 2_4 4e 0.1205 0.3344 0.0264 O 22 4e 0.121 0.3423 0.0759
O 3 4e 0.5536 0.0288 0.2763 O 14 4e 0.5128 0.9604 0.2644
O 3_2 4e 0.2763 0.4712 0.5536 O 6 4e 0.2728 0.4806 0.5407
O 3_3 4e 0.9464 0.9712 0.2237 O 4 4e 0.9456 0.0026 0.22
O 3_4 4e 0.2237 0.5288 0.9464 O 17 4e 0.1926 0.4597 0.9504
O 4 4e 0.5024 0.627 0.0674 O 12 4e 0.501 0.6373 0.0895
O 4_2 4e 0.0674 0.873 0.5024 O 11 4e 0.0674 0.8818 0.503
O 4_3 4e 0.9976 0.373 0.4326 O 10 4e 0.0247 0.3088 0.4124
O 4_4 4e 0.4326 0.127 0.9976 O 20 4e 0.4061 0.1889 6.00E-04
O 5 4e 0.6715 0.9302 0.1845 O 16 4e 0.6374 0.9432 0.1701
O 5_2 4e 0.1845 0.5698 0.6715 O 5 4e 0.1443 0.5747 0.6218
O 5_3 4e 0.8285 0.0698 0.3155 O 2 4e 0.8285 0.0925 0.3167
O 5_4 4e 0.3155 0.4302 0.8285 O 19 4e 0.3116 0.4339 0.8497
O 6 4e 0.57395 0.7267 0.25215 O 21 4e 0.6087 0.6938 0.2655
O 6_2 4e 0.25215 0.7733 0.57395 O 8 4e 0.2464 0.7854 0.5621
O 6_3 4e 0.92605 0.2733 0.24785 O 3 4e 0.9051 0.3063 0.2331
O 6_4 4e 0.24785 0.2267 0.92605 O 24 4e 0.2257 0.1829 0.8819

40

Table 2.6. Displacement vectors for pristine Fe2(MoO4)3.

Fe 4 -0.0217 0.0158 -0.0071 0.3628
Fe 3 -0.0023 -0.0066 -0.0037 0.0863
Fe 2 -0.0176 0.0079 -0.0081 0.2824
Fe 1 -0.0019 -0.0021 -0.0007 0.035
Mo 1 -0.0059 -0.0034 -0.013 0.2043
Mo 6 -0.0013 -0.0054 0.0014 0.0614
Mo 4 0.0152 0.0114 0.0069 0.2585
Mo 3 0.0237 -0.0078 0.0202 0.4255
Mo 2 0.0073 -0.0236 0.0037 0.2482
Mo 5 0.0152 0.0262 0.0091 0.3455
O 1 0.0406 -0.0022 -0.0057 0.6825
O 2 -0.0222 -0.0084 -0.0021 0.3565
O 3 0.0043 0.006 0.0299 0.4655
O 4 -0.0109 -0.0465 -0.0083 0.4722
O 5 0.0299 -0.0192 -0.016 0.6323
O 6 -0.0344 0.0557 -0.0199 0.7554
O 7 0.0264 0.0528 0.0387 0.8059
O 8 -0.0005 -0.0079 -0.0495 0.7934
O 9 0.0408 0.0684 0.0119 0.8926
O 10 0.0035 -0.0094 0.0129 0.2158
O 11 0.0008 -0.0314 0.0037 0.2987
O 12 0.0311 0.0691 -0.004 0.829
O 13 0.0014 -0.0103 -0.0221 0.3736
O 14 0 -0.0088 -0.0006 0.0827
O 15 -0.0271 0.0642 0.0202 0.8577
O 16 0.0265 -0.0619 -0.003 0.727
O 17 0.0341 -0.013 0.0144 0.5393
O 18 0.0402 -0.0049 0.0497 0.8626
O 19 0 -0.0227 -0.0012 0.2129
O 20 0.0039 -0.0037 -0.0212 0.3645
O 21 -0.0348 0.0329 -0.0134 0.6153
O 22 0.0057 -0.0121 0.0118 0.2164
O 23 0.0209 -0.033 0.0147 0.4647
O 24 0.0221 0.0438 0.0442 0.802

41

Figure 2.11. Amplimodes analysis between pristine Fe2(MoO4)3 and lithiated
Li2Fe2(MoO4)3, with positions of Fe, Mo, and O in pristine Fe2(MoO4)3 and position of Li
in Li2Fe2(MoO4)3 converted into the same reference structure setting. The transformation
vectors are plotted on each atom in (a) and filtered by amplitude of 0.4 Å in (b) for clarity.
FeO6 octahedra are shown in green, and MoO4 tetrahedra are shown in gray. Li
+
ions are
left out for amplimodes analysis but are drawn in the reference structure for clarity.

The most prominent displacement seen from these vectors are those associated with
oxygen atoms. The vectors on Fe and Mo atoms are extremely small, indicating that the
center of mass for the rigid polyhedral subunits remains fixed as they rotate to enlarge the
unit cell volume. It is interesting to note, as shown in Figure 2.11b, that Li
+
ions are inserted
between two neighboring lantern units and, as a result, there are significant displacement
of the oxygen atoms toward the guest ions, presumably a consequence of local bonding or
electrostatic interactions. This suggests that the ability for the framework to distort in such

42

a way that Li–O bonds can form and break in a reversible fashion is an important
consideration in the design of efficient intercalation hosts.
2.3.2. Na
+
insertion

Figure 2.12a shows the galvanostatic cycling of Fe2(MoO4)3 against Na/Na
+
. An initial
capacity of 79 mA·h·g
−1
was found, which corresponds to the insertion of 1.7 Na
+
per
Fe2(MoO4)3, which is comparably smaller than that observed with Li
+
insertion. The
capacity was found to drop to 87% of the initial capacity after 20 cycles, showing a similar
capacity retention to that of the Li/Fe2(MoO4)3 cell. The most pronounced difference from
Li
+
insertion is that there are two distinct slope regions during the Na
+
-intercalation process
instead of a single plateau. This is signified by two peaks (centered at ca. 2.7 V and ca.
2.58 V) in the derivative curve given as the inset of Figure 2.12a. The reproducibility of
these peaks with respect to cell potential during subsequent cycles demonstrates that the
Na
+
de(insertion) process is also highly reversible, and the electrochemical profile is
intrinsic to the Na
+
-intercalation process.
The in situ XRD experiments during Na
+
intercalation showed substantial differences
compared to that of Li
+
insertion. During discharge, a constant shift in the Bragg reflections
toward lower angles was observed with all the peaks returning to their original positions
after fully charging. This shift corresponds to the expansion of the unit cell volume in order
to accommodate the guest alkali cations during the insertion process, but given that no new
reflections evolve during cycling, this indicates that a solid solution mechanism is at play,
in which the active material stoichiometry changes continuously, in contrast to the biphasic
process described before for lithium.

43

Figure 2.12. (a) Galvanostatic electrochemical cycling of Fe2(MoO4)3 against Na
+

insertion and its derivative (shown as inset). There are two slope regions during both
insertion and de-insertion. The turn from one region to the other corresponds to 0.8 Na
+

per formula being (de)inserted. (b) 2D pattern based on the in situ XRD of Na
+
insertion
into Fe2(MoO4)3.

While the laboratory XRD patterns did not provide enough intensity to allow for a full
refinement of the atomic positions, it was possible to track changes in the unit cell using
Le Bail fits to the inserted patterns. Relative changes in these parameters are plotted in
Figure 2.13 with a complete list of the refined lattice parameters given in Table 2.7. A
volume change of 6% was found for insertion of 1.5 Na
+
ions per formula unit, with the
dominant change in the unit cell corresponding to a roughly 4% elongation of the b-axis.
While the magnitude of the change along the a- and c-axes is less pronounced, the trend

44

seems to suggest the underlying insertion mechanism. The c-axis does not significantly
change during the initial stages of Na
+
insertion, while the a-axis exhibits a continuous
extension. After the insertion of 0.8 Na
+
per formula unit, the trends are reversed (i.e., the
change in a-axis flattens out while the c-axis shows a continuous increase). It is also
important to recognize that the unit cell parameters of the Na
+
intercalated monoclinic
phase and the orthorhombic structure identified for Li2Fe2(MoO4)3 suggests that the parent
framework of Fe2(MoO4)3 actually rotates in the same way regardless of Li
+
or Na
+

insertion. The primary difference distinction between the two space group choices depends
heavily on the position where the guest ions end up sitting. Within this context, the
pronounced elongation of the b-axis is probably the result of the larger Na
+
ions relative to
that of Li
+
and that the size of the intercalant plays a crucial role in determining how the
structure can distort.

Figure 2.13. Lattice evolution along Na
+
insertion into Fe2(MoO4)3 from Le Bail fitting of
in situ X-ray diffraction patterns.

45

Table 2.7. Lattice parameters of NaxFe2(MoO4)3 from Le Bail fit of in situ XRD patterns.
x values are calculated based on number of electrons flown into the cell.

x a (Å) b (Å) c (Å) β (deg) Volume (Å
3
)
0 15.7405 9.2353 15.7011 109.1237 2156
0.1 15.77 9.25 15.69 109.09 2163
0.2 15.7749 9.2784 15.7031 109.01 2172
0.3 15.8064 9.3178 15.6971 108.97 2186
0.4 15.853 9.332 15.69 108.8 2197
0.5 15.858 9.35 15.72 108.72 2208
0.6 15.875 9.372 15.72 108.56 2216
0.7 15.885 9.393 15.691 108.31 2223
0.8 15.894 9.425 15.696 108.19 2234
0.9 15.887 9.47 15.736 108.18 2249
1.0 15.893 9.525 15.771 108.12 2269
1.1 15.889 9.564 15.804 108.12 2283
1.2 15.899 9.587 15.845 107.96 2297
1.3 15.906 9.614 15.883 107.83 2312
1.4 15.902 9.626 15.888 107.87 2315
1.5 15.938 9.614 15.901 107.93 2318
Li2Fe2(MoO4)3 in P21/c setting
16.007 9.338 16.007 107.33 2284

Given the complexity of the monoclinic structure, which contained 34 unique atomic
positions and over 150 refinable parameters, it proved impossible to directly refine the
structure of Na
+
-inserted Fe2(MoO4)3 by starting from the model of the pristine structure.
Therefore, we turned to density functional theory calculations to obtain a predicted
structure for monoclinic Na2Fe2(MoO4)3, which was used for further Rietveld refinement.
We first calculated an optimized structure for orthorhombic Li2Fe2(MoO4)3, starting from
the experimentally determined geometry. This gave excellent agreement with the Rietveld
refined geometry (e.g., lattice parameters agree to within 0.5%), which demonstrates that

46

our PBEsol+U calculations are able to well describe the structure for this high-symmetry
lithium-ordered phase. The atomic positions from the DFT-optimized orthorhombic
Li2Fe2(MoO4)3 were then projected onto a monoclinic cell as a starting structure for the
Na2Fe2(MoO4)3 structure optimization. The calculated Li
+
-inserted and Na
+
-inserted
structures from the simulations are included in Tables 2.8 and 2.9. Despite the extremely
good agreement between the DFT cell for the lithiated compound and its experimentally
determined structure, there is a striking difference between the increase of the unit cell
volume for Na
+
-insertion predicted by DFT (16.0%) and that found experimentally (6.0%).
Notwithstanding this discrepancy, the DFT-optimized Na2Fe2(MoO4)3 structure generates
a close match for the peak position and intensities in the neutron diffraction data (Figure
2.14). Using the DFT-optimized structure for subsequent Rietveld refinement, however,
produces unphysically large atomic displacement parameters for the sodium atoms. These
large displacement parameters indicate that the intensity contributed to the pattern by Na
+

is negligible and suggests a disordered sodium configuration in the experimental sample.
Refining the optimized structure again after removing sodium

atoms produces a fit nearly
identical to the refinement with the sodium included (Figure 2.15). The resulting
framework is illustrated in Figure 2.16 with the full structural description listed in Table
2.10.

47

Table 2.8. Unit cell parameters and atomic positions of optimized Li-inserted Fe2(MoO4)3
from PBEsol+U DFT calculations.

Space Group Pbcn
a (Å) 12.8839
b (Å) 9.5242
c (Å) 9.3350
Unit cell volume (Å
3
) 1145.49

Atom x y z Occupancy Wyckoff
Mo1 0.000006 0.464404 0.250005 1.000 4c
Mo2 0.357247 0.404085 0.101795 1.000 8d
Fe1 0.376885 0.247000 0.465556 1.000 8d
O3 0.163103 0.091825 0.106497 1.000 8d
O6 0.094892 0.352982 0.166014 1.000 8d
O5 0.220677 0.171506 0.462356 1.000 8d
O4 0.064418 0.571549 0.377238 1.000 8d
O2 0.012296 0.860786 0.066933 1.000 8d
O1 0.179065 0.170553 0.770329 1.000 8d
Li 0.190137 0.223019 0.265684 1.000 8d

Table 2.9. Unit cell parameters and atomic positions of optimized Na
+
-inserted
Fe2(MoO4)3 from PBEsol+U DFT calculations.

Space Group P1
a (Å) 16.57179
b (Å) 19.23659
c (Å) 9.63158
α (deg) 89.99951
β (deg) 89.99951
γ (deg) 54.52504
Unit cell volume (Å
3
) 2500.440

48

______________________________________________
Atom x y z Occup. Wyckoff
Mo1 0.000001 0.236452 0.249982 1.000 1a
Mo2 0.000001 0.736452 0.249982 1.000 1a
Mo3 0.999988 0.263551 0.750010 1.000 1a
Mo4 0.999988 0.763551 0.750010 1.000 1a
Mo5 0.500000 0.763545 0.749980 1.000 1a
Mo6 0.500000 0.263545 0.749980 1.000 1a
Mo7 0.500003 0.236454 0.250011 1.000 1a
Mo8 0.500003 0.736454 0.250011 1.000 1a
Mo9 0.357362 0.020783 0.109866 1.000 1a
Mo10 0.357362 0.520783 0.109866 1.000 1a
Mo11 0.642638 0.979216 0.890133 1.000 1a
Mo12 0.642638 0.479216 0.890133 1.000 1a
Mo13 0.142635 0.979219 0.609866 1.000 1a
Mo14 0.142635 0.479219 0.609866 1.000 1a
Mo15 0.857361 0.020782 0.390128 1.000 1a
Mo16 0.857361 0.520782 0.390128 1.000 1a
Mo17 0.642700 0.878137 0.390062 1.000 1a
Mo18 0.642700 0.378137 0.390062 1.000 1a
Mo19 0.357302 0.121861 0.609938 1.000 1a
Mo20 0.357302 0.621861 0.609938 1.000 1a
Mo21 0.857293 0.621870 0.890070 1.000 1a
Mo22 0.857293 0.121870 0.890070 1.000 1a
Mo23 0.142701 0.378131 0.109923 1.000 1a
Mo24 0.142701 0.878131 0.109923 1.000 1a
Fe1 0.385760 0.931257 0.470291 1.000 1a
Fe2 0.385760 0.431257 0.470291 1.000 1a
Fe3 0.614233 0.068752 0.529714 1.000 1a
Fe4 0.614233 0.568752 0.529714 1.000 1a
Fe5 0.114238 0.068739 0.970293 1.000 1a
Fe6 0.114238 0.568739 0.970293 1.000 1a
Fe7 0.885768 0.931241 0.029719 1.000 1a
Fe8 0.885768 0.431241 0.029719 1.000 1a
Fe9 0.614253 0.817014 0.029693 1.000 1a
Fe10 0.614253 0.317014 0.029693 1.000 1a
Fe11 0.385738 0.182983 0.970315 1.000 1a
Fe12 0.385738 0.682983 0.970314 1.000 1a
Fe13 0.885745 0.682994 0.529697 1.000 1a
Fe14 0.885745 0.182994 0.529697 1.000 1a
Fe15 0.114261 0.317013 0.470309 1.000 1a
Fe16 0.114261 0.817013 0.470309 1.000 1a
O1 0.149346 0.968203 0.098293 1.000 1a
O2 0.149346 0.468203 0.098293 1.000 1a

49

O3 0.850646 0.031799 0.901697 1.000 1a
O4 0.850646 0.531799 0.901697 1.000 1a
O5 0.350652 0.031792 0.598295 1.000 1a
O6 0.350652 0.531792 0.598295 1.000 1a
O7 0.649345 0.968208 0.401704 1.000 1a
O8 0.649345 0.468208 0.401704 1.000 1a
O9 0.850610 0.617543 0.401711 1.000 1a
O10 0.850610 0.117543 0.401711 1.000 1a
O11 0.149390 0.382455 0.598283 1.000 1a
O12 0.149390 0.882455 0.598283 1.000 1a
O13 0.649388 0.882455 0.901706 1.000 1a
O14 0.649388 0.382455 0.901706 1.000 1a
O15 0.350605 0.117548 0.098289 1.000 1a
O16 0.350605 0.617548 0.098289 1.000 1a
O17 0.081245 0.139895 0.154178 1.000 1a
O18 0.081245 0.639895 0.154178 1.000 1a
O19 0.918753 0.860102 0.845837 1.000 1a
O20 0.918753 0.360102 0.845837 1.000 1a
O21 0.418748 0.860105 0.654177 1.000 1a
O22 0.418748 0.360105 0.654177 1.000 1a
O23 0.581240 0.139898 0.345829 1.000 1a
O24 0.581240 0.639898 0.345829 1.000 1a
O25 0.918785 0.721126 0.345860 1.000 1a
O26 0.918785 0.221126 0.345860 1.000 1a
O27 0.081217 0.278873 0.654162 1.000 1a
O28 0.081217 0.778873 0.654162 1.000 1a
O29 0.581219 0.778870 0.845857 1.000 1a
O30 0.581219 0.278870 0.845857 1.000 1a
O31 0.418772 0.221136 0.154151 1.000 1a
O32 0.418772 0.721136 0.154151 1.000 1a
O33 0.265010 0.028419 0.999795 1.000 1a
O34 0.265010 0.528419 0.999795 1.000 1a
O35 0.734989 0.971578 0.000211 1.000 1a
O36 0.734989 0.471578 0.000211 1.000 1a
O37 0.234984 0.971584 0.499797 1.000 1a
O38 0.234984 0.471584 0.499797 1.000 1a
O39 0.765014 0.028418 0.500202 1.000 1a
O40 0.765014 0.528418 0.500202 1.000 1a
O41 0.735021 0.793445 0.500182 1.000 1a
O42 0.735021 0.293445 0.500182 1.000 1a
O43 0.264978 0.206556 0.499825 1.000 1a
O44 0.264978 0.706556 0.499825 1.000 1a
O45 0.764976 0.706558 0.000193 1.000 1a
O46 0.764976 0.206558 0.000193 1.000 1a

50

O47 0.235021 0.293443 0.999808 1.000 1a
O48 0.235021 0.793443 0.999808 1.000 1a
O49 0.426138 0.825448 0.362476 1.000 1a
O50 0.426138 0.325448 0.362476 1.000 1a
O51 0.573854 0.174549 0.637513 1.000 1a
O52 0.573854 0.674549 0.637513 1.000 1a
O53 0.073858 0.174552 0.862465 1.000 1a
O54 0.073858 0.674552 0.862465 1.000 1a
O55 0.926149 0.825454 0.137514 1.000 1a
O56 0.926149 0.325454 0.137514 1.000 1a
O57 0.573891 0.751583 0.137592 1.000 1a
O58 0.573891 0.251583 0.137592 1.000 1a
O59 0.426104 0.248426 0.862399 1.000 1a
O60 0.426104 0.748426 0.862399 1.000 1a
O61 0.926104 0.748419 0.637581 1.000 1a
O62 0.926104 0.248419 0.637581 1.000 1a
O63 0.073899 0.251565 0.362402 1.000 1a
O64 0.073899 0.751565 0.362402 1.000 1a
O65 0.474427 0.931991 0.054540 1.000 1a
O66 0.474427 0.431991 0.054540 1.000 1a
O67 0.525572 0.068009 0.945465 1.000 1a
O68 0.525572 0.568009 0.945465 1.000 1a
O69 0.025565 0.068014 0.554547 1.000 1a
O70 0.025565 0.568014 0.554547 1.000 1a
O71 0.974434 0.931987 0.445453 1.000 1a
O72 0.974434 0.431987 0.445453 1.000 1a
O73 0.525644 0.906336 0.445300 1.000 1a
O74 0.525644 0.406336 0.445300 1.000 1a
O75 0.474356 0.093662 0.554701 1.000 1a
O76 0.474356 0.593662 0.554701 1.000 1a
O77 0.974358 0.593659 0.945306 1.000 1a
O78 0.974358 0.093659 0.945306 1.000 1a
O79 0.025642 0.406335 0.054689 1.000 1a
O80 0.025642 0.906335 0.054689 1.000 1a
O81 0.334054 0.006389 0.285716 1.000 1a
O82 0.334054 0.506389 0.285716 1.000 1a
O83 0.665941 0.993614 0.714283 1.000 1a
O84 0.665941 0.493614 0.714283 1.000 1a
O85 0.165946 0.993607 0.785718 1.000 1a
O86 0.165946 0.493607 0.785718 1.000 1a
O87 0.834056 0.006391 0.214273 1.000 1a
O88 0.834056 0.506391 0.214273 1.000 1a
O89 0.666053 0.840425 0.214217 1.000 1a
O90 0.666053 0.340425 0.214217 1.000 1a

51

O91 0.333945 0.159573 0.785782 1.000 1a
O92 0.333945 0.659573 0.785782 1.000 1a
O93 0.833938 0.659588 0.714228 1.000 1a
O94 0.833938 0.159588 0.714228 1.000 1a
O95 0.166059 0.340412 0.285764 1.000 1a
O96 0.166059 0.840412 0.285764 1.000 1a
Na1 0.187724 0.018555 0.280487 1.000 1a
Na2 0.187724 0.518555 0.280487 1.000 1a
Na3 0.812279 0.981438 0.719502 1.000 1a
Na4 0.812279 0.481438 0.719502 1.000 1a
Na5 0.312278 0.981440 0.780497 1.000 1a
Na6 0.312278 0.481440 0.780497 1.000 1a
Na7 0.687719 0.018563 0.219501 1.000 1a
Na8 0.687719 0.518563 0.219501 1.000 1a
Na9 0.812309 0.706235 0.219481 1.000 1a
Na10 0.812309 0.206235 0.219481 1.000 1a
Na11 0.187695 0.293765 0.780510 1.000 1a
Na12 0.187695 0.793765 0.780510 1.000 1a
Na13 0.687682 0.793773 0.719468 1.000 1a
Na14 0.687682 0.293773 0.719468 1.000 1a
Na15 0.312314 0.206226 0.280530 1.000 1a
Na16 0.312314 0.706226 0.280530 1.000 1a

Figure 2.14. Rietveld refinement against neutron diffraction data of the chemically Na
+
-
inserted Fe2(MoO4)3 collected at the POWGEN beamline at the Oak Ridge National
Laboratory using the calculated Na2Fe2(MoO4)3 model listed in Table S9 (χ
2
= 6.393, Rwp
= 4.43%).

52

Figure 2.15. Rietveld refinement of the Fe2(MoO4)3 framework derived from the
calculated Na2Fe2(MoO4)3 (χ
2
= 6.399, Rwp = 4.43%).

Figure 2.16. Resulting Fe2(MoO4)3 framework from Rietveld refinement against neutron
diffraction data using optimized Na2Fe2(MoO4)3 from simulations with sodium omitted as
the starting model.

53

Table 2.10. Resulting unit cell parameters and atomic positions of the Fe2(MoO4)3
framework from Ritveld refinement agasint the neutron diffraction data of Na
+
-inserted
Fe2(MoO4)3 using calculated Na2Fe2(MoO4)3 with Na
+
ions omitted.

Space Group P2 1/c
a (Å) 9.48901 (29)
b (Å) 9.34670 (23)
c (Å) 16.00705 (23)
β (deg) 126.35610 (20)
Unit cell volume (Å
3
) 1143.336793
_______________________________________________________________________
Atom x y z Uiso Occ. Wyck.
Mo1 0.459410(5) 0.750000 0.000000 0.028800(13) 1.000 4e
Mo2 0.042740(5) 0.886940(9) 0.639840(4) 0.028800(13) 1.000 4e
Mo3 0.236940(6) 0.113060(5) 0.139840(7) 0.028800(13) 1.000 4e
Fe1 0.872040(4) 0.529750(5) 0.622530(4) 0.018800(15) 1.000 4e
Fe2 0.373020(5) 0.470250(5) 0.122530(4) 0.018800(15) 1.000 4e
O1 0.920300(5) 0.891500(7) 0.836400(7) 0.020000(8) 1.000 4e
O2 0.752500(5) 0.108500(7) 0.336400(7) 0.020000(8) 1.000 4e
O3 0.260500(5) 0.836300(7) 0.909400(7) 0.020000(8) 1.000 4e
O4 0.558300(5) 0.163700(7) 0.409400(7) 0.020000(8) 1.000 4e
O5 0.050000(5) 0.039300(7) 0.720300(7) 0.020000(8) 1.000 4e
O6 0.390600(5) -0.039300(7) 0.220300(7) 0.020000(8) 1.000 4e
O7 0.632200(5) 0.624400(7) 0.565200(7) 0.020000(8) 1.000 4e
O8 0.498200(5) 0.375600(7) 0.065200(7) 0.020000(8) 1.000 4e
O9 0.871600(5) 0.933200(7) 0.513900(7) 0.020000(8) 1.000 4e
O10 0.156200(5) 0.066800(7) 0.013900(7) 0.020000(8) 1.000 4e
O11 0.003500(5) 0.737200(7) 0.678700(7) 0.020000(8) 1.000 4e
O12 0.353900(5) 0.262800(7) 0.178700(7) 0.020000(8) 1.000 4e

The lack of scattered intensity associated with the sodium atom positions suggests a
high degree of disorder for these ions throughout the Fe2(MoO4)3. Considered alongside
the erroneous predicted volume increase for the DFT-optimized structure, these results
suggest that the calculated volume expansion is a consequence of an (artificial) ordering of
Na
+
within the polyanion framework. The use of a (1 × 2 ×1) periodic cell for the DFT
calculations imposes an artificial symmetry on the Na
+
ions, and this may block a fully
disordered sodium configuration, instead predicting an ordered Na
+
-inserted structure. The
larger size of Na
+
versus Li
+
, however, means this Na
+
ordering can only be accommodated
within the polyanion framework through a significant volume expansion. This conceptual

54

model predicts that the smaller volume expansion and apparent Na
+
disorder observed in
the experimental refined structure are coupled. We propose that in the experimental system
the lattice strain produced by intercalating large Na
+
ions are accommodated by disordering
of these ions, rather than the otherwise necessary volume expansion predicted in the higher
symmetry DFT calculations. In contrast, smaller Li
+
ions are able to occupy an ordered
arrangement of sites upon intercalation, without imposing a large volume change, leading
to highly consistent DFT-predicted and experimental structures. The inability of our DFT
calculations to predict structures in close agreement with those obtained from experiment
for Na
+
-inserted Fe2(MoO4)3 indicates the limitations of standard periodic zero-
temperature calculations for modeling disordered intercalation phases, and suggests
alternative approaches that consider thermally disordered systems, such as large-scale
molecular dynamics, may be necessary to obtain a detailed atomic-scale description of
these disordered ions.

Figure 2.17. Experimental total scattering data for pristine, Li
+
-inserted, and Na
+
-inserted
Fe2(MoO4)3.

55

To further characterize the nature of the insertion process on the local structure of
Fe2(MoO4)3, total neutron scattering data was collected on the pristine and chemically
inserted samples (Figure 2.17). The peak observed at 1.76 Å, which corresponds to the
Mo–O interatomic distance within the MoO4 tetrahedra, does not change at all after Li
+
/Na
+

insertion. This is consistent with the more covalent and therefore more rigid nature of the
MoO4 tetrahedra. In contrast, there is a very significant elongation of the Fe–O bonds that
should be expected as the oxidation state of Fe changes from 3+ to 2+ upon insertion of the
alkali ions. It is also interesting to point out that that the long-range correlations, especially
those at long r values > 3 Å, are significantly more dampened in the Na
+
inserted than that
of the pristine or Li
+
-inserted phases. This reflects the local disorder that results as a
consequence of the interactions with the disordered distribution of Na
+
through the lattice.
Finally, Raman spectra of Li
+
- and Na
+
-inserted Fe2(MoO4)3 were collected to
determine if signatures of the disorder could be identified (Figure 2.4b). Both the lithiated
and sodiated samples show Raman bands belonging to the MoO4 units that appear to be
either weak or unresolved. The strongest asymmetric stretching band of the MoO4 units in
the Raman spectrum of the pristine material appears to be very broad and weak in the
spectra of lithiated and sodiated samples. However, the weaker bending and symmetric
stretching modes remained unresolved. In addition, no shift in the positions of Raman
bands was observed. This broadening and loss of intensity may be attributed to the
redistribution of electron density in the Mo–O bonds through the intercalation of alkali
guest ions into the pristine material, which may also be affected by the reduction of Fe
3+
ions to Fe
2+
.
24,25
Therefore, the effective force constants and polarizability derivatives may
be varied as the electron density is redistributed, resulting in the observed broadening and

56

intensity loss in the Raman bands. The major difference was observed in the asymmetric
stretching band of MoO4 tetrahedra. The intensity of the strongest asymmetric stretching
mode at 776 cm
–1
was found to be lower than of the second asymmetric stretching mode at
817 cm
–1
in the intercalated samples. This suggests that strong interactions between alkali
ions and the polyhedral oxygen ions results in changed intensities for stretching motions in
the MoO4 units, while the framework remains largely the same.

Figure 2.18. Illustration of a representative unit of the refined framework of Li
+
- and Na
+
-
inserted Fe2(MoO4)3. Li
+
from Rietveld refinement of Li2Fe2(MoO4)3 is illustrated for
perspective.

57

Table 2.11. Dihedral angles of O-Mo-Fe-O in refined frameworks of the Li
+
- and Na
+
-
inserted Fe2(MoO4)3. Dihedral angle values based on the Fe and Mo on the top-left corner
in Figure 2.18 are included.

Angle Na-inserted Li-inserted
O2-Mo2-(O1)-Fe1-O2 43.15 42.59
O2-Mo2-(O1)-Fe1-O3 125.09 123.89
O2-Mo2-(O1)-Fe1-O4 49.9 51.18
O2-Mo2-(O1)-Fe1-O5 141.29 144.91
O2-Mo2-(O1)-Fe1-O6 8.98 12.93
O3-Mo2-(O1)-Fe1-O2 75.51 72.59
O3-Mo2-(O1)-Fe1-O3 6.43 8.71
O3-Mo2-(O1)-Fe1-O4 168.56 166.36
O3-Mo2-(O1)-Fe1-O5 100.05 99.91
O3-Mo2-(O1)-Fe1-O6 127.64 128.11
O5-Mo2-(O1)-Fe1-O2 152.93 155.15
O5-Mo2-(O1)-Fe1-O3 125.13 123.55
O5-Mo2-(O1)-Fe1-O4 59.88 61.38
O5-Mo2-(O1)-Fe1-O5 31.51 32.35
O5-Mo2-(O1)-Fe1-O6 100.8 99.63

With a description of the Li
+
- and Na
+
-inserted phases of Fe2(MoO4)3, a comparison
between the changes to the structural frameworks should provide insight into the
mechanism for intercalation. A representative portion of the refined framework with
lithium determined by Li2Fe2(MoO4)3 refinement is illustrated in Figure 2.18. To
understand how the polyhedral rotations are affected by the alkali ions, a set of dihedral
angles are calculated from the structural models. The center atoms of two connected
polyhedra form two planes with one oxygen atom from each of the polyhedra in these
dihedral angles, which thus measures the relative rotational displacement between the two
polyhedra. For example, the tweaking between the FeO6 octahedra and MoO4 tetrahedra

58

can be measured by the dihedral angles of such as O2-Mo2-Fe1-O5. A complete list of the
dihedral angles between these two polyhedra is given in Table 2.11. It can be seen there is
a difference in every dihedral angle between the Na
+
-inserted and Li
+
-inserted Fe2(MoO4)3,
and a difference smaller than 2° can be considered as systematic residuals. There are more
significant differences, such as the 3° difference in the O2-Mo2-Fe1-O5 and O2-Mo2-Fe1-
O6 angles. From our symmetry mode analysis, the transformation from the pristine to the
inserted phase is accompanied by oxygen atoms moving towards the inserted ion (Li
+
),
possibly driven by a strong electrostatic Li–O interaction. Thus, as O5 on the bottom left
corner in Figure 2.18 moves towards the alkali ion, driving the FeO6 to rotate in a way that
essentially results in O5 and O6 moving away from O2 in the MoO4 tetrahedra on the top
left corner, and because the Li–O bond is shorter than the Na–O bond, the dihedral angle
of O2–Mo2–Fe1–O5 is larger and that of O2–Mo2–Fe1–O6 is smaller in Li
+
-inserted
Fe2(MoO4)3 than Na
+
-inserted phase. This again highlights that while a striking difference
has been observed in the electrochemical cycling curves, as well as structural phase
changes between the Li
+
- and Na
+
-intercalation in Fe2(MoO4)3, the fundamental
mechanism for Li
+
- and Na
+
-intercalation is effectively the same but structural differences
arrise because of the differences in ionic radii.
We have investigated the insertion mechanisms of two different alkali guest ions into
the Fe2(MoO4)3 framework and found that the ability for the framework to distort via
cooperative rotations appears to be critical for facilitating intercalation. This has been
confirmed by the symmetry-mode analysis of pristine and Li
+
-inserted Fe2(MoO4)3, and
through careful analysis of the total scattering data on both Li
+
- and Na
+
-inserted samples.
Similar observations have been found in LiFeSO4F cathode materials,
26–28
and the lack of

59

connection between these two structure types implies that such a mechanism may have
wider applications in understanding polyanionic electrode materials. It should also be noted
that the distortion through rigid and strongly covalently bonded polyhedra restricts the way
the host structure can distort, and thus offers better longevity than oxide materials, which
usually suffer from drop in performance due to severe structural changes during
cycling.
29,30

While the mechanism of polyhedral rotational distortion we have proposed has not been
widely discussed with respect to intercalation electrodes, it has been extensively explored
in the superionic conductor literature.
31–33
This is typically found in studies of two
competing mechanisms: (i) a “paddle wheel” mechanism, which suggests a strong
correlation between the transport of cations and the rotation of polyhedral subunits, and (ii)
a “percolation” mechanism, which claims independence between conduction pathway and
polyhedral rotations.
31,32
For example, through powder neutron diffraction and reverse
Monte Carlo (RMC) modeling, Karlsson et al. revealed that both types of mechanisms are
present in high-temperature forms of Li2SO4 and LiNaSO4.
31,33

These rotational distortions clearly suggest a generalized mechanism of alkali ion
intercalation in polyanionic materials, but more importantly, this work has revealed a
significant contrast in the behavior between Li
+
- and Na
+
-intercalation into the same host,
which suggest that differences in the intrinsic nature of Li
+
and Na
+
guest ions may result
in a modification of such a mechanism. We observed a two-stage, solid solution Na
+
insertion process as compared to a single, two-phase Li
+
insertion process, which is similar
to what has been found in olivine-type FePO4 and explained by stronger interactions of Na
+

ions as compared to Li
+
ions with the host structure by Moreau et al.
34
Our results also

60

indicate that the two redox peaks in Fe2(MoO4)3 do not appear to correspond to the
complete filling of two distinct Na
+
positions, but rather that Na
+
ions adopt a disordered
distribution throughout the lattice. This disorder is likely the result of two discrete
crystallographic positions, each having a distinct chemical potential for the alkali ion, but
as the framework rotates and distorts to accommodate more ions the Na
+
ions already in
the framework are pushed off their ideal position. A more careful investigation into
intermediate compositions may provide more valuable information in this regard, but a
comparison between the total neutron scattering data on Li
+
- and Na
+
-inserted Fe2(MoO4)3
shows that peaks at high r values are less intense in Na
+
-inserted sample than those in Li
+
-
inserted sample, which reinforces the notion that the rotational distortions are less coherent
following Na
+
than Li
+
insertion. Both observations can be explained by the larger ionic
size of Na
+
than Li
+
, and thus a steric hindrance for the rotational distortions to allow the
Na
+
ions to find an ideal position. Finally, we note that the highly insulating character of
these materials would require that the charge transport through the lattice would have to be
mediated through some kind of polaronic hopping of the electrons. It could be possible that
the structural distortions reported here are related to this mechanism of charge hopping,
and further work using X-ray absorption spectroscopy may be useful in elucidating this
aspect of the transport properties.

61

2.4. Experimental

2.4.1. Synthesis of Fe2(MoO4)3
Fe2(MoO4)3 was synthesized by a solid-state precipitation method.
16

(NH4)6Mo7O24·4H2O (1.0 mmol) and Fe(NO3)3·9H2O (4.65 mmol) were dissolved
separately in 100 mL and 50 mL of H2O, respectively. 0.4 mL of NH3·H2O was added to
the molybdate solution to establish basic conditions. The Fe(NO3)3 solution was then
slowly added to the molybdate solution with stirring, resulting in a yellow precipitate. The
reaction mixture was stirred overnight. The resulting yellow precipitate was collected and
washed with ethanol by sonicating for 30 min, and isolated by centrifugation (6000 rpm
for 15 min). This washing procedure was repeated twice. The resulting Fe2(MoO4)3 powder
was dried in a vacuum oven at 60
˚
C for 5 h. The dried Fe2(MoO4)3 was then ground with
a mortar and pestle and annealed at 400
˚
C for 6 h in air. A yellowish-green powder was
obtained.
2.4.2. Chemical insertion
Li2Fe2(MoO4)3 was prepared by stirring 1.0 g Fe2(MoO4)3 in fivefold excess of LiI
(dissolved in dry acetonitrile) for 2 weeks. For Na
+
insertion, a sodium napthalenide
solution was prepared by dissolving 3.0 g of Na metal and 1.5 g of naphthalene in 75 mL
of dry THF. Na2Fe2(MoO4)3 was prepared by mixing 1.0 g Fe2(MoO4)3 in 25 mL of the
resulting sodium napthalenide solution for a week.

62

2.4.3. Characterization
Laboratory powder X-ray diffraction (XRD) data was collected on a Bruker D8
diffractometer with a CoKα source (λ1 = 1.78897 Å, λ2 = 1.79285 Å), equipped with a
LynxEye detector. The collection process was kept the same for different samples using a
0.6 mm slit with a step size of 0.02˚ and a total collection time of 1 h in the 2θ range from
10 to 60 degrees.
A Swagelok-type electrochemical cell with a beryllium disk as current collector for the
working electrode was used for in situ XRD as well as for diffraction data collection over
chemically lithiated and sodiated samples.
35,36
Laboratory XRD data were all collected in
reflectance.
High-resolution synchrotron powder diffraction data were collected at room
temperature using 11-BM beamline at the Advanced Photon Source (APS), Argonne
National Laboratory using an average wavelength of 0.413965 Å. The data were collected
in the 2θ range of -6.5° to 28.0° with a step size of 0.001˚ and 0.1 second spent on each
step. Neutron diffraction and total scattering data were collected at room temperature on
polycrystalline powders loaded in a vanadium can using the time-of-flight POWGEN
instrument (BL-11A) and NOMAD instrument (BL-1B), respectively, at the Spallation
Neutron Source, Oak Ridge National Laboratory. POWGEN data were collected with 16
hours spent on banks 3 and 4 hours on bank 4. Bank 3 uses frame #1.5 at 60 Hz with a Q
range of 1.1688 Å
-1
to 15.1700 Å
-1
and Bank 4 uses frame #1.75 at 60 Hz with a Q range
of 0.6820 Å
-
1 to 5.6923 Å
-1
. Data from all detectors with angles from 20° to 150° were
combined into one histogram for each bank. NOMAD data were collected in the scattering
angle range of 3° to 175° with 40 mins spent on each sample. Data from all detectors were

63

reduced into one PDF pattern after background contribution was subtracted. The resulting
data were analyzed using the General Structure Analysis System (GSAS)
37
and PDFgui.
38

The high-resolution X-ray photoelectron spectroscopy (XPS) spectrum was acquired
using a Kratos Axis Ultra X-ray photoelectron spectrometer with the analyzer lens in
hybrid mode. A monochromatic aluminum anode with an operating current of 6 mA and
voltage of 10 kV was used with a step size of 0.1 eV, a pass energy of 20 eV, and a pressure
range between 1  10
–8
torr to 3  10
–8
torr. The binding energy was referenced to the C1s
core level at 284 eV.
Raman spectra were recorded in the 200–1200 cm
–1
range using a Horiba Xplora
Raman microscope (Horiba Scientific). Laser irradiation of 532 nm wavelength was
employed as the excitation source and the power at the sample level was 50 mW. All
spectra were recorded under ambient conditions, a quartz cell was used for the
measurement of lithiated and sodiated samples. Intercalated samples were loaded into the
quartz cell in an Ar-filled glove box.
2.4.4. Electrochemical measurements
Electrochemical measurements were carried out using Swagelok-type cells.
36
The
positive electrodes were prepared by grinding Fe2(MoO4)3 and carbon Super P (CSP) for
30 min in an Ar-filled glove box. A slurry was prepared by adding the resulting mixture
into polyvinylidenefluoride (PVDF) dissolved in N-methyl-2-pyrrolidone (NMP). The
slurry, comprised of 76.4 wt% Fe2(MoO4)3, 8.5 wt% CSP, and 15.1 wt% PVDF, was casted
onto an aluminum foil and dried in a vacuum oven at 110
˚
C overnight. Electrodes were
peeled off the Al foil and punched into 1.0 cm diameter circles. Li or Na metal disks were
used as the negative electrodes. Whatman GF/D borosilicate glass fiber sheets were used

64

as the separator and soaked in electrolyte solution of either 1 M LiPF6 in ethylene carbonate
(EC), propylene carbonate (PC), and dimethylcarbonate (DMC) (45:45:10 by weight) or 1
M NaClO4 in EC and DMC (1:1 by weight).
2.4.5. Symmetry-mode analysis
Symmetry-mode analysis was carried out using the AMPLIMODES program of the
Bilbao Crystallographic Server.
23
The inputs of the program include structural parameters
of a high-symmetry and a low-symmetry structure. Results from Rietveld refinements
against synchrotron X-ray and neutron diffraction data on pristine and Li
+
-inserted
Fe2(MoO4)3 were used as low-symmetry and high-symmetry phases, respectively.
2.4.6. Computational
Density functional theory (DFT) calculations were performed using the plane-wave
code VASP.
39,40
Interactions between core and valence electrons were described with the
PAW method, with cores of [Kr] for Mo, [Ar] for Fe, [He] for O, and [Ne] for Na.
41
All
calculations used the PBEsol exchange–correlation functional.
42
To describe the strongly
correlated Fe 3d electrons, we applied a Hubbard-type “+U” correction of Ud = 4.3 eV,
using the rotationally invariant approach of Dudarev et al.
43
This value of Ud = 4.3 eV was
chosen for consistency with previous calculations of orthorhombic Li
+
- and Na
+
-inserted
Fe2(MoO4)3 by Yue et al.
15
who, in turn, selected this value from earlier calculations on
LixFePO4.
44
A planewave cutoff of 550 eV was used, and all calculations were spin-
polarized. High spin Fe was assumed, extrapolating from calculations of the parent
monoclinic Fe2(MoO4)3 phase that predicted a high-spin anti-ferromagnetic solution. All
structures were geometry-optimized with no assumed symmetry by performing a series of
constant volume structural relaxations and fitting the resultant data to the Murnaghan

65

equation of state. Individual geometry optimizations were deemed converged when all
atomic forces fell below 0.01 eV Å
-1
. Orthorhombic Li2Fe2(MoO4)3 was modelled using a
single unit cell (76 atoms), with 2 × 2 × 2 Monkhorst-Pack k-point sampling, using our
Rietveld-refined structure as a starting configuration. Because we were unable to directly
refine the monoclinic Na2Fe2(MoO4)3 structure, an approximate structure used to initialize
the DFT geometry optimization was constructed by projecting the optimized
Li2Fe2(MoO4)3 structure onto a (1 × 2 × 1)-expanded monoclinic cell (152 atoms). The
structural relaxation of the Na
+
-intercalated phase used a 2 × 1 × 2 Monkhorst-Pack k-point
grid.

2.5. Conclusions

In summary, we have presented a detailed structural characterization of the process by
which Li
+
and Na
+
ions intercalate into Fe2(MoO4)3, and how these structural
transformations are related to the correlated rotations of rigid polyanionc subunits. The
combination of electrochemical cycling and in situ powder XRD confirmed a two-phase
process for Li
+
-intercalation, compared to a two-stage solid solution process for Na
+
-
intercalation, which corresponds to Na
+
filling two different sites. Furthermore, based on
our Rietveld refinements, symmetry-mode analysis and examination of neutron total
scattering data, we proposed a concerted polyhedra rotational distortion mechanism for the
alkali guest ion intercalation into the anti-NASICON Fe2(MoO4)3 host framework. It is also
shown that during Li
+
-insertion, Li
+
ions fill defined positions that allow the transformation
from the pristine monoclinic phase to the lithiated orthorhombic structure. However, Na
+
insertion occurs in a disordered manner that appears to result from the larger size of Na
+

compared to Li
+
. Such a mechanism, as well as the application of symmetry-mode analysis

66

along with structural and electrochemical characterization, may be applied to other
polyanionic material systems. Insights gained through these analyses may facilitate the
discovery of new intercalation materials and the improvement of existing ones.
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70

Chapter 3. Investigating the Mechanism of Reversible Lithium Insertion into Anti-
NASICON Fe2(WO4)3*
*Published in ACS Appl. Mater. Interfaces 2017, 9, 10813-10819.

3.1. Abstract

The gram-scale preparation of anti-NASICON Fe2(WO4)3 by a new solution-based
route and detailed characterization of the material are presented. The resulting Fe2(WO4)3
undergoes a reversible electrochemical reaction against lithium centered around 3.0 V with
capacities near 93% of the theoretical maximum. Structural evolution of Fe2(WO4)3 upon
lithium (de)insertion is probed by the combination of in situ X-ray diffraction, X-ray
absorption spectroscopy (XAS), and neutron total scattering techniques, coupled with
symmetry-mode analysis. A two-phase structural transformation from monoclinic to
orthorhombic is confirmed during lithium intercalation. XAS and neutron total scattering
measurements verify that Fe2(WO4)3 consists of trivalent iron and hexavalent tungsten ions.
As lithium ions are inserted into the framework, iron ions are reduced to the divalent state,
while the tungsten ions are electrochemically inactive and remain in the hexavalent state.
Lithium insertion occurs via a concerted rotation of the rigid polyhedra in the host lattice
driven by interactions with the Li
+
ions; the magnitude of these rotational distortions was
found to be slightly larger for Fe2(WO4)3 than for the Fe2(MoO4)3 analog.
3.2. Introduction

Rechargeable Li-ion batteries have found utility in applications ranging from portable
electronics to electric vehicles.
1–3
The majority of cathodes used in commercialized Li-ion
batteries rely on layered lithium transition metal oxides (i.e., LiMO2, where M = Co, Mn,

71

and Ni).
4,5
In oxides, the distortions that occur as Li
+
ions are intercalated into the lattice
can induce severe changes in the structural framework leading to capacity fade over time.
Therefore, significant effort has been devoted to searching for alternative cathode materials
with better longevity, such as polyanion materials that have structural frameworks
composed of corner-sharing polyhedral subunits. Since the first reports by Goodenough et
al., the structural and electrochemical properties of so-called polyanion compounds with
rigid tetrahedral structural units (i.e., XO4; X = S, P, Mo, etc.) have been extensively
investigated.
6,7
These polyanion compounds exhibit better thermal stability when
compared to layered lithium transition metal oxides owing to the covalent X-O bonds in
the tetrahedral subunits. Moreover, the redox potential of polyanion compounds can be
tuned via polyanion substitution by modification of the ionic character of the M-O bond,
which is known as the inductive effect.
8,9
NASICON-type frameworks were among the
first polyanion compounds investigated as cathodes in Li-ion batteries.
10–14
These
materials, possessing the general formula MM’(XO4)3, are structurally comprised of
covalently bonded 3D frameworks. The number of inserted alkali ions per formula unit
depends on the available oxidation states of the transition metal (M) in the octahedra and
the element (X) in the tetrahedra.
9,15

Owing to the robust structural topology created by large interstitial voids scattered
throughout a three dimensionally connected network of polyhedra, anti-NASICON type
Fe2(MO4)3 (where M = Mo or W) is a promising intercalation electrode material. The
electrochemical performance of both Li
+
and Na
+
intercalation into Fe2(MoO4)3 has been
studied in some detail.
14,16–18
Recently, we reported the mechanistic difference between Li
+

and Na
+
(de)insertion into monoclinic Fe2(MoO4)3. It was found that Li
+
intercalation

72

proceeds via concerted rotation of the rigid polyhedral subunits, driven by strong
interactions with the Li
+
ions, which ultimately leads to an ordered lithium arrangement in
the anti-NASICON structure. While Na
+
intercalation occurs in a fundamentally similar
manner, sodium insertion results in a far less coherent final structure because of its
substantially larger ionic radius.
19
In contrast to Fe2(MoO4)3, the structural and
electrochemical properties of Fe2(WO4)3 remain largely unexplored. This is possibly
because of the lower theoretical specific capacity (63 mA h g
–1
) for this material owing to
the heavy mass of tungsten. However, the excellent electrochemical cycling performance
of this material, resulting from its extremely robust structural topology, makes it a
fundamentally interesting material in which to study the mechanism of lithium
intercalation.
Ferric tungstate was first reported in 1965 as a compound of unknown structure that
was purported to melt at 1065 ˚C.
20
Subsequently, Pernicone and Fagherazzi prepared a
phase-impure mixture of “Fe2W3O12” and WO3 through a precipitation method followed
by annealing at 700 ˚C under flowing oxygen, and identified the new iron tungstate as
having a tetragonal unit cell.
21
In 1985, Harrison et al. were the first to successfully prepare
true Fe2(WO4)3 and reported it to be isostructural with Fe2(MoO4)3 and crystallize in a
monoclinic (P21/a) space group.
22
Sriraman and Tyahi confirmed this assignment and
reported the thermal decomposition of Fe2(WO4)3 to FeWO4 and WO3 at temperatures
above 750 ˚C.
23
Monoclinic Fe2(WO4)3 crystallizes into a 3D structure consisting of FeO6
octahedra and WO4 tetrahedra interconnected through corner sharing oxygen atoms.
22
In
contrast to the isostructural Fe2(MoO4)3, Fe2(WO4)3 does not have a high-temperature
orthorhombic polymorph.
23
To date, the structural aspects of lithium intercalation into

73

Fe2(WO4)3 have only been investigated by Goodenough et al. through chemical insertion
with ex situ powder X-ray diffraction.
24
They showed that chemical Li
+
insertion into
Fe2(WO4)3 results in a fully lithiated Li2Fe2(WO4)3 phase that crystallizes into a structure
indexed to an orthorhombic (Pbcn) space group in the same way as Fe2(MoO4)3. Up to this
point, detailed structural characterization of anti-NASICON Fe2(WO4)3, the mechanism
through which Li
+
ions electrochemically (de)insert into the material, and how the
structural evolution is correlated with electrochemical cycling, has not been reported.
Herein, we report a new solution precipitation route for the gram-scale synthesis of phase-
pure anti-NASICON Fe2(WO4)3. We investigate the structural transformation of the
pristine material into orthorhombic Li2Fe2(WO4)3 with respect to Li
+
intercalation using a
combination of structural and electrochemical tools and demonstrate the correlation
between structural evolution and electrochemical properties of Fe2(WO4)3.

3.3. Results and Discussion

Attempts to prepare Fe2(WO4)3 through a high-temperature solid-state route were first
reported by Nassau et al., but were found to be unsuccessful.
20
Following this work, there
have been a handful of successful examples of lower-temperature synthesis of monoclinic
Fe2(WO4)3 through solution-based precipitation methods.
22,23
The difficulty in preparing
phase-pure monoclinic Fe2(WO4)3 arises from its thermal instability at high temperatures,
with traditional solid-state routes leading to the formation of the thermodynamically stable
FeWO4 phase via disproportionation.
20,23
Previous routes to monoclinic Fe2(WO4)3 using
solution precipitation methods relied on the use of a Na2WO4 precursor, which can lead to
sodium-inserted products.
22
Therefore, we developed an alternative precipitation approach

74

for the synthesis of monoclinic Fe2(WO4)3 using (NH4)10W12O41·xH2O and
Fe(NO3)3·9H2O as the starting materials according to following reaction:
8Fe(NO3)3∙9H2O + (NH4)10W12O41∙xH2O → 4Fe2(WO4)3 + 34NH3 + yH2O

Figure 3.1. XRD patterns of as-prepared powders annealed at various temperatures.

This method relies on the room temperature precipitation of an amorphous
nanoparticulate product with subsequent two-stage thermal annealing. The washed and
dried amorphous nanoparticles are first annealed at 475 ˚C for 8 h, followed by an
additional annealing step at 550 ˚C for 8 h. This graded annealing provides better
crystallinity than direct thermal aging of the amorphous product at 550 ˚C for 8 h (Figure
3.1). Direct thermal annealing at 600 ˚C or thermal annealing above 650 ˚C results in the
decomposition of Fe2(WO4)3 into FeWO4 and WO3 instead of a phase transition from the
monoclinic phase into a high temperature orthorhombic phase, as is the case for
Fe2(MoO4)3.
23
This synthetic approach results in the preparation of large quantities of
phase-pure Fe2(WO4)3, with isolated ceramic yields of 95%.

75

Figure 3.2. (a) Fit of the monoclinic model to the experimental PDF of pristine Fe2(WO4)3
neutron total scattering data obtained at the NOMAD beamline at Oak Ridge National
Laboratory and (b) the resulting crystal structure. (Iron, tungsten and oxygen are shown as
green, purple and orange, respectively.)

The phase purity and crystal structure of the pristine, monoclinic Fe2(WO4)3 was
confirmed by neutron total scattering data collected at NOMAD beamline at Oak Ridge
National Laboratory. Modeling was constrained to the (P21/c) space group and initialized
with structural parameters obtained from the model of Chen et al. for monoclinic
Fe2(MoO4)3.
25
The fit to the experimental PDF of monoclinic Fe2(WO4)3 in the 1.6–15 Å
range is shown in Figure 3.2. Detailed structural parameters including atomic positions,
cell parameters, and isotropic displacement Table 3.1. The low value of Rw, as well as visual
inspection of the fit to the experimental data, confirm that Fe2(WO4)3 is isostructural with

76

Fe2(MoO4)3 and that the structure is well-described by the monoclinic (P21/c) anti-
NASICON structure (Figure 3.2b).

Table 3.1. Resulting unit cell, atomic positions, and atomic displacement parameters from
the fit of the monoclinic model to the experimental PDF of pristine Fe2(WO4)3 neutron total
scattering data obtained at the NOMAD beamline at Oak Ridge National Laboratory.

Atom Wyckoff Occupancy x y z Uiso (Å
2
)
W1 4e 1.0 0.512600 0.746550 0.505470 0.013028
W2 4e 1.0 0.862680 0.628180 0.223590 0.000998
W3 4e 1.0 0.751020 0.612410 0.890230 0.005349
W4 4e 1.0 0.614790 0.116040 0.765830 0.013850
W5 4e 1.0 0.787890 0.130850 0.139110 0.039619
W6 4e 1.0 -0.019820 0.244580 -0.017120 0.026869
Fe1 4e 1.0 0.681010 0.465100 0.061670 0.015472
Fe2 4e 1.0 -0.044820 -0.040700 0.325220 0.005932
Fe3 4e 1.0 0.814760 -0.026200 -0.06614 0.010064
Fe4 4e 1.0 0.584080 0.483000 0.689960 0.017334
O1 4e 1.0 0.000100 0.887200 0.581800 0.019090
O2 4e 1.0 0.828500 -0.092500 0.816700 0.023906
O3 4e 1.0 -0.094900 0.693700 0.733100 0.015835
O4 4e 1.0 -0.054400 -0.002600 0.720000 0.010219
O5 4e 1.0 0.855700 -0.074700 0.378200 0.023808
O6 4e 1.0 0.727200 0.019400 0.459300 0.019374
O7 4e 1.0 0.586100 0.609900 0.006500 0.058498
O8 4e 1.0 0.753600 0.785400 -0.062100 0.027480
O9 4e 1.0 0.556300 0.853600 0.104500 0.033871
O10 4e 1.0 0.024700 0.808800 0.412400 0.006341
O11 4e 1.0 -0.067400 0.881800 -0.003000 0.000390
Space Group P21/c
a (Å) 15.91510
b (Å) 9.30921
c (Å) 15.61460
β (
o
) 108.95600

77

O12 4e 1.0 0.499000 0.862700 -0.089500 0.008798
O13 4e 1.0 0.767500 0.882900 0.625000 0.094814
O14 4e 1.0 0.487200 0.539600 0.735600 0.011230
O15 4e 1.0 0.667600 0.606500 0.78400 0.006523
O16 4e 1.0 0.637400 0.443200 0.170100 0.007939
O17 4e 1.0 0.807400 0.459700 0.549600 0.007150
O18 4e 1.0 0.699300 0.437800 0.364700 0.009527
O19 4e 1.0 0.688400 0.433900 0.650300 0.025988
O20 4e 1.0 0.406100 0.811100 0.500600 0.013380
O21 4e 1.0 0.608700 0.306200 0.765500 0.008557
O22 4e 1.0 0.879000 0.157700 -0.075900 0.002780
O23 4e 1.0 0.877200 0.102700 0.240100 0.012763
O24 4e 1.0 0.774300 0.317100 0.118100 0.037077

Figure 3.3. High-resolution XPS spectra of Fe (left) and W (right) in pristine Fe2(WO4)3.

The chemical composition of the pristine Fe2(WO4)3 was determined by ICP-OES to
possess a Fe/W ratio of 0.65, in close agreement with the nominal composition of the
material. X-ray photoelectron spectroscopy (XPS) gave Fe 2p3/2 and W 4f7/2 binding
energies of 711.3 and 35.6 eV, respectively, with spin-orbit splitting matching closely with
literature values for Fe
3+
and W
6+
.
26,27
Additionally, a broad peak at 41.5 eV corresponds
to the expected W 5p binding energy for W
6+
(Figure 3.3). Transmission electron

78

microscopy (TEM) was used to gain insight into the morphology of the resulting
monoclinic Fe2(WO4)3 (Figure 3.4). TEM images before and after annealing to 550 ˚C
reveal that thermal treatment leads to the aggregation and grain growth of the amorphous
nanoparticles, with an increase in primary particle size from ~50 to ~200 nm being
observed.

Figure 3.4. TEM images of Fe2(WO4)3 (a) before and (b) after annealing to 550 ˚C.

The electrochemical behavior of monoclinic Fe2(WO4)3 toward Li
+
(de)insertion was
investigated by galvanostatic cycling in concert with in situ X-ray diffraction. The first ten
galvanostatic electrochemical cycles of Fe2(WO4)3 against Li/Li
+
at a cycling rate of C/10
are given in Figure 3.5a. It was determined that 93% of the theoretical capacity of two Li
+

ions per formula unit was achieved on the first discharge. The derivative of the
galvanostatic cycling gives two peaks centered at 3.01 and 2.96 V during oxidation and
reduction, respectively (Figure 3.5b). Such a small polarization suggests an excellent
reversibility of lithium (de)insertion into the anti-NASICON framework operating on the
Fe
3+
/Fe
2+
redox couple. Without rigorous optimization of the cell assembly, the cells
displayed a high capacity retention of around 95% of the initial capacity after 25 cycles at
a galvanostatic cycling rate of C/10. The voltage-composition curve exhibits a single
plateau over a wide range of Li content during Li
+
(de)insertion, which suggests that

79

intercalation occurs through a two-phase process where a lithium-rich phase is formed
directly in increasing mass fraction with increasing lithium content, rather than through a
continuous solid-solution process.
24,28

Figure 3.5. (a) Galvanostatic electrochemical cycling of Fe2(WO4)3 against Li
+
insertion
and (b) its derivative curve.

To confirm that lithium (de)insertion is occurring through a two-phase process, the
structural evolution of monoclinic Fe2(WO4)3 with lithium (de)insertion was monitored by
in situ X-ray diffraction. Figure 3.6 displays a heat map of in situ XRD patterns being taken
continuously during electrochemical cycling by a rate of C/10 with the y-axis
corresponding to inserted and extracted lithium content upon discharging and charging,
respectively. The reflections of pristine, monoclinic Fe2(WO4)3 progressively vanish

80

without any shift in peak positions. As an illustration, Figure 3.6 shows that the intensities
of reflections at 25.1˚, 26.6˚, 31.8˚, 34.9˚, and 39.5˚ gradually decrease until finally
disappearing as the Li
+
content increases during discharging. This observation is
characteristic of a two-phase insertion process, with a new set of diffraction peaks (i.e., at
24.3˚, 25.6˚, 25.8˚, 26.3˚, 29.4˚, 31.0˚, 38.3˚, and 42.5˚) appearing upon discharging since
the lithiated phase crystallizes into a different structure. This new lithium-inserted
Li2Fe2(WO4)3 phase can be indexed to an orthorhombic structure with a Pbcn space group,
as originally posited by Goodenough and coworkers.
24
During the charging process,
changes in the intensities of reflections were observed in the opposite way indicating a
highly reversible structural variation upon Li
+
(de)insertion, as suggested by
electrochemical cycling data (Figure 3.5). These variations in the intensities of reflections
are illustrative of the evolution from the monoclinic pristine structure to an orthorhombic
lithiated phase upon intercalation. The coexistence of Bragg reflections associated with the
both pristine and lithiated phases confirms the two-phase mechanism for the phase
transformation.

81

Figure 3.6. 2D pattern of in situ XRD study for a full discharge/charge cycle of Fe2(WO4)3
against lithium at C/10 rate with a CoKα source (λ1 = 1.78897 Å, λ2 = 1.79285 Å). Color
indicates the intensities of reflections.

While the structure of Li2Fe2(WO4)3 was previously postulated to be isostructural to
Li2Fe2(MoO4)3,
24
we sought to verify this by chemically inserting Li
+
ions into Fe2(WO4)3
by soaking powders in a solution of LiI in dry acetonitrile so that high-quality neutron total
scattering data could be collected. The final composition of the lithiated phase was
determined by ICP-OES to have a Li/Fe/W ratio of 0.72:0.66:1. The fit of the experimental
PDF on the chemically lithiated phase in the 1.6–15 Å range is given in Figure 3.7a. The
experimental PDF of lithiated phase was modeled to the Pbcn space group using the
structural parameters starting from the model of Manthiram et al. for Li2Fe2(WO4)3.
24,29
A
high quality-of-fit was achieved with a low R wp value of 15%. Detailed structural
parameters including cell parameters, atomic positions, and isotropic atomic displacement
parameters are given in Table 3.2. The excellent fit of the PDF extracted from neutron total
scattering data to the orthorhombic phase confirms that the fully lithiated anti-NASICON
frameworks of Li2Fe2(WO4)3 and Li2Fe2(MoO4)3 are indeed isostructural (Figure 3.7b).

82

Table 3.2. Resulting unit cell, atomic positions, and atomic displacement parameters from
the fit of the orthorhombic model to the experimental PDF of lithiated Li2Fe2(WO4)3
neutron total scattering data obtained at the NOMAD beamline at Oak Ridge National
Laboratory.

Space Group Pbcn
a (Å) 12.79510
b (Å) 9.52608
c (Å) 9.36247
Atom Wyckoff Occupancy x y z Uiso (Å
2
)
W1 8d 1.0 0.140180 0.092860 0.101240 0.034304
W2 4c 1.0 0.000000 0.534750 0.250000 0.002714
Fe1 8d 1.0 0.124450 0.258410 0.467640 0.016149
O1 8d 1.0 0.013260 0.148870 0.065750 0.011745
O2 8d 1.0 0.055500 0.430970 0.376020 0.012696
O3 8d 1.0 0.179170 0.171440 0.269350 0.009759
O4 8d 1.0 0.273190 0.339150 0.464580 0.006326
O5 8d 1.0 0.341010 0.410330 0.111590 0.017876
O6 8d 1.0 0.396230 0.138180 0.167240 0.008679
Li1 8d 1.0 0.321840 0.287240 0.278110 0.003070

83

Figure 3.7. (a) Fit of the orthorhombic model to the experimental PDF of lithium-inserted
Li2Fe2(WO4)3 neutron total scattering data obtained at the NOMAD beamline at Oak Ridge
National Laboratory and (b) resulting crystal structure. (Iron, tungsten, oxygen and lithium
are shown as green, purple, orange and yellow, respectively.)

Given that tungsten can adopt both a hexavalent and tetravalent oxidation state, we
sought to interrogate whether the redox center associated with Li (de)insertion in
Fe2(WO4)3 is fully isolated to the iron ions in the structure. In this regard, X-ray absorption
near edge structure (XANES) spectroscopy has proven extremely powerful for determining
the oxidation state of metals in polyanion cathode materials.
30,31
XANES measurements
were performed on Fe2(WO4)3 before and after lithium intercalation. The Fe K-edge
XANES spectra of pristine and lithiated samples compared against FeSO 4 are given in
Figure 3.8a. The position of the absorption edge in pristine Fe2(WO4)3 (i.e., 7124.9 eV) is
shifted toward lower energy upon lithium intercalation, with the absorption edges of

84

Li2Fe2(WO4)3 and FeSO4 (i.e., 7119.2 and 7119.1 eV, respectively) being nearly identical
and conforming to the change in iron oxidation state from 3+ to 2+ upon lithiation.
31

Potential variations in the electronic structure of the polyanion units before and after
lithium intercalation were probed through W LIII-edge XANES spectroscopy. The XANES
regions of W LIII-edge in Fe2(WO4)3 and Li2Fe2(WO4)3 are given in Figure 3.8b. Both
pristine and lithiated samples exhibit a characteristic sharp asymmetrical peak for
tetrahedral W units, which are attributed to electron transitions from 2p3/2 to 5d states.
These W LIII absorption edges of the pristine and lithiated phases are identical, indicating
that lithium intercalation does not affect the electronic structure around tungsten.

Figure 3.8. Normalized (a) Fe K-edge and (b) W LIII-edge XANES spectra, and (c)
experimental neutron total scattering data for Fe2(WO4)3 and chemically lithiated
Li2Fe2(WO4)3.

The neutron total scattering data for the pristine and chemically lithiated Fe2(WO4)3
were compared to probe the nature of the intercalation process on the local structure of the
anti-NASICON framework. Experimental pair distribution functions of pristine and
lithiated phases in the range of 1.0–2.6 Å are given in Figure 3.8c. The PDF peak at 1.77
Å corresponds to the W–O interatomic distance within the tungstate tetrahedra and does

85

not change with lithium intercalation, consistent with the rigid nature of these polyhedral
subunits. Conversely, the Fe–O interatomic distance increases from 2.00 to 2.09 Å after
lithium insertion into the framework, consistent with a change in oxidation state for Fe
from 3+ to 2+ on lithium insertion and in agreement with the XANES data. Collectively,
this confirms that lithium (de)insertion operates only on Fe
3+
/Fe
2+
redox couple.

Figure 3.9. The transformation vectors on each atom from the symmetry-mode analysis
between pristine Fe2(WO4)3 and lithiated Li2Fe2(WO4) as distorted phase and parent phase,
respectively. Positions of Fe, W, and O atoms in Fe2(WO4)3 and position of lithium in
Li2Fe2(WO4)3 converted into the same reference structure setting (P21/c). Li
+
ions are left
out for the symmetry-mode analysis but are drawn in the reference structure for clarity.
(Iron, tungsten, oxygen, and lithium atoms are shown as green, purple, orange, and yellow,
respectively.)

To more rigorously describe the structural distortion of the monoclinic Fe2(WO4)3
phase into orthorhombic phase, symmetry-mode analysis by the AMPLIMODES
program at the Bilbao Crystallographic Server was performed.
32
The transformation of
the low-symmetry pristine structure into the higher-symmetry lithiated one was
identified by using a transformation matrix (Table 3.3). The higher-symmetry structure
(Pbcn) was converted into the low-symmetry (P21/c) one after subtraction of lithium

86

atoms from the lithiated structure, and then atoms in these two structures were directly
paired (Table 3.4 and Table 3.5). The displacement vectors on each atom for the
structural transformation were calculated and illustrated in Figure 3.9, and the complete
list of precise vectors is given in the Table 3.6. The dominant distortions observed from
the symmetry-mode analysis are displacements of the oxygen atoms. The vectors on Fe
and W atoms are negligible, indicating the preservation of center of mass of the
polyhedral subunits during rotational distortions upon Li
+
insertion. In other words, the
structural transformation during lithium intercalation into Fe2(WO4)3 is driven by a
rotational displacement of oxygen atoms toward the Li
+
ions, which likely stems from
the creation of new electrostatic interactions with the framework.

Table 3.3. Transformation matrix for converting parent structure (Pbcn) into reference
structure (P21/c) for symmetry mode analysis.

[ 1 0 -1 ] [ 1 ]
[ 1 0 1 ] [ 1/2 ]
[ 0 -1 0 ] [ 0 ]

Table 3.4. Unit cell parameters of parent structure into reference structure settings.

Space Group P21/c
a (Å) 15.951827
b (Å) 9.362470
c (Å) 15.951827
β (
o
) 106.663895

87

Table 3.5. Atomic pairing for the parent phase and the distorted phase in the reference
settings.

Reference Structure Low Symmetry Structure
Atom Wyckoff x y z Atom Wyckoff x y z
O1 4e 0.33107 0.93425 0.31781 O 16 4e 0.3626 0.9432 0.3299
O1_2 4e 0.31781 0.56575 0.33107 O 19 4e 0.3116 0.5661 0.3497
O1_3 4e 0.16894 0.06575 0.1822 O 2 4e 0.1715 0.0925 0.1833
O1_4 4e 0.1822 0.43425 0.16894 O 5 4e 0.1443 0.4253 0.1218
O2 4e 0.49323 0.62398 0.43773 O 12 4e 0.499 0.6373 0.4105
O2_2 4e 0.43773 0.87602 0.49323 O 20 4e 0.4061 0.8111 0.5006
O2_3 4e 0.00677 0.37602 0.06227 O 10 4e 0.9753 0.3088 0.0876
O2_4 4e 0.06227 0.12398 0.00677 O 11 4e 0.0674 0.1182 0.003
Fe1 4e 0.44143 0.53236 0.31698 Fe 4 4e 0.41592 0.517 0.31004
Fe1_2 4e 0.31698 0.96764 0.44143 Fe 1 4e 0.31899 0.9651 0.43833
Fe1_3 4e 0.05857 0.46764 0.18302 Fe 2 4e 0.04482 0.4593 0.17478
Fe1_4 4e 0.18302 0.03236 0.05857 Fe 3 4e 0.18524 0.0262 0.06614
W1 4e 0.36652 0.89876 0.22634 W 4 4e 0.38521 0.88396 0.23417
W1_2 4e 0.22634 0.60124 0.36652 W 5 4e 0.21211 0.63085 0.36089
W1_3 4e 0.13348 0.10124 0.27366 W 2 4e 0.13732 0.12818 0.27641
W1_4 4e 0.27366 0.39876 0.13348 W 3 4e 0.24898 0.38759 0.10977
O3 4e 0.4253 0.73065 0.24613 O 21 4e 0.3913 0.6938 0.2345
O3_2 4e 0.24613 0.76935 0.4253 O 24 4e 0.2257 0.8171 0.3819
O3_3 4e 0.07469 0.26935 0.25387 O 3 4e 0.0949 0.3063 0.2669
O3_4 4e 0.25387 0.23065 0.07469 O 8 4e 0.2464 0.2146 0.0621
O4 4e 0.55617 0.53542 0.28298 O 14 4e 0.5128 0.4604 0.2644
O4_2 4e 0.28298 0.96458 0.55617 O 6 4e 0.2728 0.9806 0.5407
O4_3 4e 0.94383 0.46458 0.21702 O 4 4e 0.9456 0.5026 0.22
O4_4 4e 0.21702 0.03542 0.94383 O 17 4e 0.1926 0.9597 0.9504
O5 4e 0.62567 0.88841 0.28466 O 15 4e 0.6676 0.8935 0.284
O5_2 4e 0.28466 0.61159 0.62567 O 18 4e 0.3007 0.5622 0.6353
O5_3 4e 0.87433 0.11159 0.21534 O 23 4e 0.8772 0.1027 0.2401
O5_4 4e 0.21534 0.38841 0.87433 O 13 4e 0.2325 0.3829 0.875
O6 4e 0.51721 0.83276 0.12098 O 9 4e 0.5563 0.8536 0.1045
O6_2 4e 0.12098 0.66724 0.51721 O 22 4e 0.121 0.6577 0.5759
O6_3 4e 0.98279 0.16724 0.37903 O 1 4e 0.9999 0.1128 0.4182
O6_4 4e 0.37903 0.33276 0.98279 O 7 4e 0.4139 0.3901 0.9935
W2 4e 0.51738 0.75 0.51738 W 1 4e 0.5126 0.74655 0.50547

88

W2_2 4e 0.98262 0.25 0.98262 W 6 4e 0.98018 0.24458 0.98288

From this, we can draw some mechanistic comparisons between lithium (de)insertion
into anti-NASICON type Fe2(MO4)3, where M = W or Mo. A discharge voltage of 2.96 V
for Fe2(WO4)3 is slightly lower than that of Fe2(MoO4)3 (3.0 V), but is so close that the
differences are most likely due to polarization from differences in the processing of the
cathode active material (i.e., efficiency of the carbon coating, particle morphology/size,
etc.).
30
The plateau in the galvanostatic voltage-composition curves for both Fe2(WO4)3
and Fe2(MoO4)3 suggest a similar two-phase structural transformation on lithium
(de)insertion, which was further corroborated by in situ X-ray diffraction experiments. As
in the case of Fe2(MoO4)3, we find reversible variations in the intensities of reflections
during the (de)insertion of Li
+
into Fe2(WO4)3, with no remarkable shifts in the positions
of the Bragg reflections being observed. Both galvanostatic electrochemical cycling and in
situ XRD experiments point towards a transformation from the monoclinic form into the
orthorhombic Li2Fe2(WO4)3 phase through a two-phase process. The structure of fully
lithiated Li2Fe2(WO4)3 was also confirmed to be orthorhombic with a Pbcn space group
through PDF analysis of neutron total scattering data on the chemically inserted material.
This lithiated phase appears to be isostructural with that of Li2Fe2(MoO4)3. Since
Fe2(WO4)3 and Fe2(MoO4)3 are isostructural for both the pristine and lithiated phases, with
an identical structural topology with regards to polyhedral connectivity, we expect an
analogous structural transformation mechanism for both materials.

89

Table 3.6. Displacement vectors for pristine Fe2(WO4)3

Atom ux Uy Uz IuI
O1 0.0315 0.0089 0.0121 0.4916
O1_2 -0.0062 0.0003 0.0186 0.3392
O1_3 0.0026 0.0267 0.0011 0.2536
O1_4 -0.0379 -0.0089 -0.0471 0.8229
O2 0.0058 0.0133 -0.0272 0.4855
O2_2 -0.0316 -0.0649 0.0074 0.8197
O2_3 -0.0315 -0.0672 0.0253 0.9632
O2_4 0.0051 -0.0058 -0.0038 0.1268
Fe1 -0.0255 -0.0154 -0.0069 0.4156
Fe1_2 0.002 -0.0025 -0.0031 0.0703
Fe1_3 -0.0138 -0.0083 -0.0082 0.2344
Fe1_4 0.0022 -0.0062 0.0076 0.1293
W1 0.0187 -0.0148 0.0078 0.3199
W1_2 -0.0142 0.0296 -0.0056 0.3532
W1_3 0.0038 0.0269 0.0027 0.2603
W1_4 -0.0247 -0.0112 -0.0237 0.4728
O3 -0.034 -0.0369 -0.0116 0.6245
O3_2 -0.0204 0.0478 -0.0434 0.81
O3_3 0.0202 0.0369 0.013 0.4779
O3_4 -0.0075 -0.016 -0.0126 0.2518
O4 -0.0434 -0.075 -0.0186 0.9707
O4_2 -0.0102 0.016 -0.0155 0.2946
O4_3 0.0018 0.038 0.003 0.3592
O4_4 -0.0244 -0.0757 0.0066 0.8299
O5 0.0419 0.0051 -0.0007 0.6736
O5_2 0.016 -0.0494 0.0096 0.5295
O5_3 0.0029 -0.0089 0.0248 0.3933
O5_4 0.0172 -0.0055 0.0007 0.2757
O6 0.0391 0.0208 -0.0165 0.7682
O6_2 0.000 -0.0095 0.0587 0.9404
O6_3 0.0171 -0.0544 0.0392 0.7918
O6_4 0.0349 0.0573 0.0107 0.7566
W2 -0.0048 -0.0034 -0.0119 0.186
W2_2 -0.0024 -0.0054 0.0003 0.0648

Total distortion amplitude (global distortion): 6.5051 Å

90

We recently demonstrated that Fe2(MoO4)3 undergoes a transformation into the
orthorhombic phase via the cooperative rotation of rigid tetrahedral subunits as Li
+
is
inserted.
19
More specifically, the molybdate tetrahedra maintained a rigid Mo–O bond
length, but were found to pivot around their crystallographic position so that the oxygen
ions could displace toward the new position of the Li
+
ions. The symmetry-mode analysis
of the transformations are nearly identical for both Fe2(WO4)3 and Fe2(MoO4)3, as might be
expected from the similar ionic radii of 4-coordinate, hexavalent Mo and W atoms (0.42 Å
and 0.41 Å, respectively);
33
however, the magnitude of the total distortion was found to be
slightly larger for Fe2(WO4)3. The amplitude of the global distortion, which represents a
sum of the displacement for all atoms in the unit cell, for Fe2(WO4)3 and Fe2(MoO4)3 were
determined to be 6.50 Å and 6.17 Å, respectively. This difference most likely arises from
differences in the boding within the two tetrahedral subunits. Tungsten has a Pauling
electronegativity of 2.36, compared to molybdenum, which is 2.16.
34
This suggests that the
tungstate tetrahedra should be more covalent, and through inductive effects should increase
the ionic character of the Fe–O bonds. As a result, it should be easier to distort the Fe–O
bonds away from ideal, facilitating a more pronounced magnitude in the rotational
distortion as observed.

3.4. Experimental

3.4.1. Synthesis
Fe2(WO4)3 was synthesized by a solution precipitation method under ambient
conditions. Ammonium paratungstate hydrate [(NH4)10W12O41·xH2O, 0.58 mmol] from
Alfa Aesar and iron(III) nitrate hydrate (Fe(NO3)3·9H2O, 4.65 mmol) from Sigma Aldrich

91

were dissolved separately in 150 mL and 50 mL of deionized water, respectively. NH3·H2O
(0.5 mL) was added to the paratungstate solution to establish basic conditions. The
Fe(NO3)3 solution was then slowly added to the paratungstate solution with stirring,
resulting in an orange precipitate. The reaction mixture was stirred overnight. The resulting
orange precipitate was filtered and subsequently washed with ethanol by sonication for 15
min, and then isolated by centrifugation (6000 rpm for 15 min). This washing procedure
was repeated twice. The resulting Fe2(WO4)3 powder was dried in a vacuum oven at 60

˚C
for 5 h. The dried Fe2(WO4)3 was then ground with a mortar and pestle and annealed at 475
˚C for 8 h in air. A brownish-orange powder was obtained. This powder was then annealed
at 550 ˚C for 8 h, resulting in a crystalline, light orange powder in 95% isolated ceramic
yield.
3.4.2. Chemical Insertion
Chemically lithiated Li2Fe2(WO4)3 was prepared by stirring 1.0 g of Fe2(WO4)3 in ten-
fold molar excess of LiI (dissolved in dry acetonitrile) for 2 weeks in an Ar-filled glove
box. The resulting light brown powder was washed with dry acetone in the glove box by
sonication for 15 min, and isolated by centrifugation (6000 rpm for 15 min). This washing
procedure was repeated twice. The resulting powder was dried under vacuum overnight at
room temperature.
3.4.3. Characterization
XRD patterns were collected on a Bruker D8 diffractometer with a CoKα source (λ1 =
1.78897 Å, λ2 = 1.79285 Å), equipped with a LynxEye detector. The data collection was
kept the same for different samples using a 0.6 mm slit with a step size of 0.02˚ and a total
collection time of 1 h in the 2θ range from 10 to 60˚. A Swagelok-type electrochemical cell

92

with a beryllium disk as current collector for the working electrode was used for in situ
XRD.
35
All laboratory XRD data were all collected in reflectance mode.
Transmission Electron Microscopy (TEM) images were obtained using a JEOL
JEM2100F (JEOL Ltd.) electron microscope operating at 200 kV. Samples for TEM
studies were prepared by drop casting a stable suspension of particles in ethanol on a 200
mesh Cu grid coated with a lacey carbon film (Ted Pella, Inc.).
X-ray Absorption Near Edge Structure (XANES) spectra were collected at the 20-BM-
B line of the Advanced Photon Source at Argonne National Laboratory. Fe K-edge and W
LIII-edge XAS was collected in transmission mode using gas ionization chambers filled
with N2 to monitor the incident and transmitted intensities. Samples were prepared by
spreading a thin, uniform layer of powder on Kapton tape. The incident X-ray beam size
was 1 mm - 6 mm (unfocused). XANES spectra were collected at room temperature under
flowing helium. Each spectrum was normalized by subtracting the pre-edge and applying
an edge-jump normalization using the Athena software.
36

Neutron total scattering data were collected at room temperature on polycrystalline
powders loaded in a vanadium can using the Nanoscale-Ordered Materials Diffractometer
(NOMAD, BL-1B) instrument at the Spallation Neutron Source, Oak Ridge National
Laboratory. NOMAD data were collected in the scattering angle range of 3−175˚ with 40
min spent on each sample. Custom software developed for the NOMAD instrument was
used to produce a reduced scattering function, S(Q), in which contributions from the sample
holder and vanadium can were subtracted. The Pair Distribution Function (PDF), G(r), was
obtained by direct Fourier transformation of S(Q) with a Qmax of ~32 Å
-1
. The resulting
PDFs were analyzed in PDFgui.
37

93

Symmetry-mode analysis was carried out using the AMPLIMODES program of the
Bilbao Crystallographic Server.
32
The input of the program includes structural parameters
of a high-symmetry lithiated and a low-symmetry pristine structure. Results from PDF
analysis of neutron total scattering data on pristine and chemically Li-inserted Fe2(WO4)3
were used as low-symmetry and high-symmetry phases, respectively. For direct
comparison, lithium atoms in the Li-inserted Fe2(WO4)3 were left out in this analysis.
High-resolution X-ray Photoelectron Spectroscopy (XPS) data were acquired using a
Kratos Axis Ultra X-ray photoelectron spectrometer with the analyzer lens in hybrid mode.
A monochromatic aluminum anode with an operating current of 6 mA and voltage of 10
kV was used with a step size of 0.1 eV, a pass energy of 20 eV, and a pressure range around
1-3  10
–8
torr. The binding energy was referenced to the C 1s core level at 284 eV. Samples
were prepared by pressing ~75 mg of dry powder into 13-mm diameter pellets using ~5
metric tons of pressure for 3 min and applying the pellets onto conductive carbon tape.
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. Inductively coupled plasma optical emission
spectroscopy (ICP-OES) was performed by Galbraith Laboratories, Inc. for Fe, W and Li
analyses.
3.4.4. Electrochemical Measurements
Electrochemical measurements were carried out using Swagelok-type cells.
35
The
positive electrodes were prepared by grinding Fe2(WO4)3 and Ketjen black carbon (2:1
w/w) with a mortar and pestle for 30 min in an Ar-filled glove box. Cells were assembled
in an Ar-filled glove box with a Li-metal disk as the negative electrode. Two Whatman

94

GF/D borosilicate glass fiber sheets were used as the separator. 1 M LiPF6 in ethylene
carbonate and dimethylene carbonate (1:1 w/w) solution was used as the electrolyte. Cells
were prepared by approximately 10 mg of active material grinded with Ketjen black carbon
and typically cycled at 25 ˚C between 2.4 and 3.8 V versus Li at a rate of 1.0 Li
+
per formula
unit over 10 h (C/10). Galvanostatic cycling was carried out on a BioLogic VMP3
potentiostat.

3.5. Conclusions

In summary, we presented a new solution-based precipitation method for the gram-
scale synthesis of anti-NASICON monoclinic Fe2(WO4)3 along with detailed structural
characterization. Using a combination of structural and electrochemical probes, we have
demonstrated the reversible (de)insertion of lithium into the Fe2(WO4)3 framework.
Galvanostatic cycling indicates that Fe2(WO4)3 undergoes a reversible electrochemical
reaction centered around ~3.0 V with a high capacity retention of 95% over 25 cycles at a
C/10 cycling rate. XANES and neutron total scattering studies confirm that the charge
compensation upon Li
+
(de)insertion is achieved through the Fe
3+
/Fe
2+
redox couple only.
A two-phase lithium intercalation process was confirmed by in situ X-ray diffraction
studies, and the transformation of the initial monoclinic structure into an orthorhombic
lithiated structure was demonstrated by symmetry-mode analysis to proceed through the
filling of defined crystallographic positions by inserted lithium atoms via cooperative
rotations polyhedral subunits. This comprehensive study provides constructive information
on the electrochemical performance and structural evolution of NASICON-type materials,

95

which may facilitate further development of insertion cathodes for rechargeable Li-ion
batteries.

3.6. References

(1) Etacheri, V.; Marom, R.; Elazari, R.; Salitra, G.; Aurbach, D. Challenges in the
Development of Advanced Li-ion Batteries: A review. Energy Environ. Sci. 2011,
4, 3243–3262.

(2) Lu, L.; Han, X.; Li, J.; Hua, J.; Ouyang, M. A Review on the Key Issues for Lithium-
ion Battery Management in Electric Vehicles. J. Power Sources 2013, 226, 272–
288.

(3) Andre, D.; Kim, S.-J.; Lamp, P.; Lux, S. F.; Maglia, F.; Paschos, O.; Stiaszny, B.
Future Generations of Cathode Materials: an Automotive Industry Perspective. J.
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(4) Mizushima, K.; Jones, P. C.; Wiseman, P. J.; Goodenough, J. B. Li xCoO2 (0< x<-
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Chapter 4. Phase Control in the Colloidal Synthesis of Well-Defined Nickel Sulfide
Nanocrystals *

*Published in Nanoscale 2018, 10, 16298-16306.

4.1. Abstract

Morphologically well-defined colloidal nanocrystals of Ni3S4, NiS, Ni9S8, and Ni3S2
were independently prepared through a solution-based method using N,N’-disubstituted
thioureas as the sulfur precursor. Simultaneous control of phase, composition, and
morphology of the resulting colloidal nickel sulfide nanocrystals was achieved by primarily
adjusting the reactivity of substituted thioureas as well as tuning the key reaction
parameters of temperature, precursor ratio, coordinating solvent. 1-dodecanethiol was
identified as an important reactivity-directing agent, and in the presence of 1-
dodecanethiol, sulfur-deficient Ni9S8 and Ni3S2 nanocrystals with spherical, rod-like and
multipod morphologies were prepared. In the absence of 1-dodecanethiol, more sulfur-rich
Ni3S4, α-NiS, and β-NiS nanocrystals were prepared individually. A phase progression
from metastable Ni3S4 to the α-NiS and thermodynamically preferred β-NiS phases upon
an increase in reaction temperature and S:Ni ratio was observed. This study presents, for
the first time, a systematic evaluation of factors that simultaneously control the phase and
morphology of nickel sulfide nanocrystals.
4.2. Introduction

Binary nickel sulfides exist in a multitude of compositions, including NiS2, Ni3S4, NiS,
Ni9S8, Ni7S6, Ni4S3 and Ni3S2. In 1962, Kullerud and Yund constructed a comprehensive
binary Ni–S phase diagram for nickel sulfides.
1
According to their phase diagram, Ni3S2
and Ni3S4 are low-temperature phases, and rhombohedral Ni3S2 (heazlewoodite) is the most

100

nickel-rich phase. Above 843 K, Ni3S2 transforms into a cubic polymorph.
1
Sulfur-rich
Ni3S4 (polydymite) is a metastable phase with a cubic spinel structure that
disproportionates into NiS and NiS1.03 phases above 373 K.
2,3
Stochiometric NiS has two
polymorphs –– β-NiS (millerite) is the low-temperature polymorph with a rhombohedral
structure, and hexagonal α-NiS is the high-temperature polymorph. The β-NiS to α-NiS
phase transition takes place above 650 K, and the metastable α-NiS phase has been shown
to display a metal to insulator transition.
4
Distinct phases of these binary nickel sulfides
have been utilized as electrocatalysts for hydrogen evolution reactions (HER)
5–7
and
oxygen evolution reactions (OER)
8–10
for water splitting, electrode materials for lithium-
ion
11–13
and sodium-ion batteries,
14,15
counter electrodes for dye-sensitized solar cells,
16

hydrodesulfurization catalysts,
17
and electroactive materials for supercapacitors.
18

The preparation of colloidal semiconductor nanocrystals with precise phase and
morphology control is crucial because the functional properties of nanocrystals depend on
the shape, size, crystal phase, and/or composition. However, the relatively complicated
nature of the Ni–S phase diagram makes the controlled solution-phase synthesis of phase-
pure colloidal nanocrystals with well-defined morphologies challenging.
1,19,20
The
synthetic methods previously reported for the preparation of nanoscale nickel sulfides
include solventless thermolytic decomposition,
21
solution-phase arrested precipitation,
22

thermal decomposition of nickel acetylacetonate in the presence of 1-dodecanethiol,
23
and
decomposition of nickel dithiocarbamates,
24–26
alkyl xanthates,
27
thiobiurets,
28
2-
mercaptobenzothiazole complexes,
29
(TMEDA)Ni(SCOC6H5)2,
30
and nickel polysulfide
complexes
31
as single source precursors in high-boiling solvents. In recent years, there have
been limited examples of the morphology- or phase-controlled synthesis of nickel sulfide

101

nanocrystals. For example, Korgel and coworkers reported the synthesis of nearly
monodispersed and well-defined β-NiS nanorods and nanopyramids via thermal
decomposition of nickel thiolate precursors in the presence of octanoic acid.
21
However,
Ni3S4 was found as an impurity phase in their as-synthesized β-NiS particles. More
recently, Liu et al. demonstrated the shape-controlled synthesis of Ni3S4 nanoparticles via
a one-pot colloidal synthesis.
32
By varying the precursor concentration and using different
nickel precursors, single-phase Ni3S4 nanocrystals with two different morphologies (i.e.,
nanoprisms and nanopyramids) were prepared.
There have also been several recent examples of the phase-controlled synthesis of
nanoscale nickel sulfides. Roffey et al. first reported the preparation of colloidal nickel
sulfides by using various nickel bis(dithiocarbamate) complexes as a single source
precursor. By changing reaction conditions, such as temperature or precursor
concentration, they were able to tune the phase of the resulting nickel sulfide
nanoparticulates between Ni3S4 and α-NiS.
24
In a similar study by Gervas et al., different
nickel complexes were utilized as single source precursors and the reaction temperature
and high-boiling solvents were shown to be the key factors for the phase control of nickel
sulfide nanoparticulates (i.e., Ni3S4 and Ni3S2).
26
Pan et al. synthesized Ni7S6, α-NiS, and
β-NiS nanoparticulates through the thermal decomposition of nickel acetylacetonate and
1-dodecanethiol in oleylamine or 1-octadecene. They observed that the Ni:S precursor ratio
and solvent choice play a decisive role in the phase control of nickel sulfide
nanostructures.
33
Although these particular studies establish the influence of various
reaction parameters on the control over nickel sulfide phase determination, they lack
sufficient control over particle morphology, with the as-synthesized nanocrystals being

102

polydispersed and possessing ill-defined shapes. To the best of our knowledge, the ability
to simultaneously exert control over phase, composition, and morphology has yet to be
demonstrated for nanocrystals in the binary Ni–S system.
Herein, we report a parametric study on the phase- and morphology-controlled
synthesis of colloidal nickel sulfide nanocrystals, and describe the role of temperature,
solvents, reagents, and S:Ni precursor ratio on phase determination and the morphology of
the reaction products. Through a systematic study of reaction parameters, we were able to
exert high degree of control over the phase and morphology of colloidal nickel sulfide
nanocrystals. Phase-pure nanocrystals of Ni3S4, α-NiS, Ni9S8, and Ni3S2 with well-defined
morphologies were independently prepared by tuning these key reaction parameters and
the growth kinetics of nanocrystals via 1-dodecanethiol. By developing a synthetic
methodology for the solution-phase synthesis of nanocrystals within the binary Ni–S phase
space, this work can be utilized to facilitate further studies of nickel sulfide nanocrystals,
along with their phase- and morphology-dependent properties and applications.

4.3. Results and Discussion

Nickel sulfide nanocrystals were synthesized via the fast addition of N,N’-disubstituted
thioureas into a hot solution of NiI2 in oleylamine. Substituted thioureas were particularly
chosen as the primary sulfur precursor because of their tunable conversion kinetics and
ability to form monodispersed metal sulfide nanoparticles with a high degree of batch-to-
batch consistency.
34
In the synthesis of sulfur-deficient phases (i.e., Ni9S8 and Ni3S2), 1-
dodecanethiol was also employed after the injection of primary sulfur precursor as an
important reactivity-directing agent. Whereas, sulfur-rich phases (i.e., Ni3S4, α-NiS and β-

103

NiS) were synthesized in the absence of 1-dodecanthiol by tuning the ratio of N,N’-
disubstituted thioureas and NiI2 precursors, reaction time and temperature (Figure 4.1).

Figure 4.1. Reaction pathways indicating the preparation of various Ni-S nanocrystals in
the presence or absence of 1-dodecanethiol as the reactivity-directing agent.

4.3.1. Sulfur-deficient Ni-S Phases
Synthesis of Ni3S2 nanocrystals. An N,N’-dibutyl thiourea solution in dibenzylamine (10
molar equivalents relative to NiI2) was quickly injected into a solution of NiI2 in oleylamine
at 180 ˚C, causing the greenish solution to turn black in about 3 min, indicating nucleation
of particles. After 5 min, 3.0 mL of 1-dodecanethiol was injected, and the reaction
temperature was maintained at 180 ˚C for 1 h (Table 4.1). The X-ray diffraction (XRD)
pattern of the resulting nanocrystals can be indexed to the rhombohedral heazlewoodite
structure of the Ni3S2 phase without any evidence of other crystalline Ni–S phases (Figure
2a). The resulting nanocrystals possess a multipod-type morphology, as observed by

104

transmission electron microscopy (TEM) analysis (Figure 4.2b,c). A high-resolution (HR-
TEM) micrograph of a single Ni3S2 multi-pod arm suggests that the arms are single crystals,
with an interplanar distance of d = 4.0 Å that corresponds to the (101) planes of the
rhombohedral heazlewoodite structure (Figure 4.2d).

Table 4.1. Summary of the reaction conditions to obtain various Ni–S nanocrystal phases
Phase Sulfur precursor S:Ni ratio Solvent/Ligand Temperature (˚C) Time
Ni3S2 N,N’-dibutyl thiourea 10.0
oleylamine &
1-dodecanethiol
180 1 h
Ni9S8 N,N’-diphenyl thiourea 3.0
oleylamine &
1-dodecanethiol
180 4 h
Ni3S4 N,N’-diphenyl thiourea 1.5 oleylamine 180 4 h
α-NiS N,N’-diphenyl thiourea 5.0 dodecylamine 180 5 min
β-NiS
*

N,N’-diphenyl thiourea 8.0 oleylamine 220 4 h
*
β-NiS phase also contains minor α-NiS impurities.

Figure 4.2. (a) XRD pattern of rhombohedral Ni3S2 nanocrystals. (b) TEM micrograph of
multipod-type Ni3S2 nanocrystals. (c-d) HR-TEM micrographs of an individual Ni3S2
nanocrystal.

105

It was found that 1-dodecanethiol used for this synthesis plays a decisive role in the
formation of heazlewoodite Ni3S2 nanocrystals. Under otherwise identical conditions, 10.0
molar equivalents of N,N’-dibutyl thiourea without the injection of 1-dodecanethiol gives
nanocrystals of the relatively more sulfur-rich orthorhombic Ni9S8 phase with rod-like
aggregates (Figure 4.3). The amount of thiol injected 5 min after N,N’-dibutyl thiourea also
played a significant role in phase formation; that is, when the amount of 1-dodecanethiol
was decreased to 0.3 mL, a mixture of Ni3S2 nanocrystals along with some Ni9S8 impurities
were formed. Whereas, when 1.5 mL of 1-dodecanethiol was injected, the resulting
nanoparticles crystallize mainly into heazlewoodite Ni3S2 phase, however, minor Ni9S8
impurities were still found (Figure 4.3a). These observations point out that by varying the
amount of 1-dodecanthiol, phase of nanocrystals can be tuned between Ni3S2 and Ni9S8,
and an increase in the amount of 1-dodecanethiol favors rhombohedral Ni3S2 phase over
orthorhombic Ni9S8 phase. Formation of more sulfur-deficient Ni3S2 phase can be
attributed to decrease in the growth kinetics of nanocrystals with increasing amount of 1-
dodecanthiol which will be discussed in more detail.

Figure 4.3. (a) Powder XRD patterns and (b) TEM micrographs of the products from the
reaction of NiI2 and N,N’-dibutyl thiourea (BuThU) in the presence of various amounts of
1-dodecanethiol (DDT) at 180 ˚C for 1 h.

106

Synthesis of Ni9S8 nanocrystals. Although we were able to obtain phase-pure Ni9S8
nanocrystals from the reaction of NiI2 and N,N’-dibutyl thiourea in the absence of 1-
dodecanthiol, the resulting nanoparticles were ill-defined and aggregated (Figure 4.3b).
Hence, for the synthesis of relatively more sulfur-rich monodispersed Ni9S8 nanocrystals,
more reactive N,N’-diphenyl thiourea was used as the primary sulfur precursor instead of
N,N’-dibutyl thiourea. In this case, NiI2 was reacted with 3.0 molar equivalents of N,N’-
diphenyl thiourea in oleylamine at 180 ˚C for 4 h (Table 4.1). As in the case of the Ni 3S2
nanocrystal synthesis, 1-dodecanethiol was also injected 5 min after the injection of
substituted thiourea to adjust the growth kinetics of nanocrystals (Figure 4.1). Figure 4.4
provides the X-ray diffraction pattern and transmission electron microscopy images of the
as-synthesized nanocrystals. TEM micrographs of the nanocrystals reveal that they are
highly uniform, quasi-spherical particles with an average diameter of 8.8  1.8 nm. The
histogram of nanoparticle size is given in Figure 4.4d, indicating a monomodal particle size
distribution. Powder XRD analysis confirms that the nanocrystals possess the
orthorhombic godlevskite structure of Ni9S8, with no evidence of other crystalline Ni–S
phases (Figure 4.4a). The high-resolution TEM image given in Figure 4.4c shows lattice
fringes of a single Ni9S8 nanocrystal that have an interplanar distance of 3.9 Å,
corresponding to the d-spacing of the (112) planes of the orthorhombic godlevskite
structure. To best of our knowledge, these nanocrystals are the smallest that have ever been
reported for orthorhombic Ni9S8 phase.

107

Figure 4.4. (a) XRD pattern of orthorhombic Ni9S8 nanocrystals. (b) TEM micrograph of
8.8-nm Ni9S8 nanocrystals. (c) HR-TEM micrograph of a single Ni9S8 nanocrystal. (d) Size
histogram showing the distribution of particle diameters for Ni9S8 (N = 355).

To elucidate the effect of 1-dodecanethiol on product phase, several control reactions
were performed in the presence and absence of thiol. Under otherwise identical conditions,
N,N’-diphenyl thiourea was used as sole sulfur precursor, with no 1-dodecanethiol being
injected subsequent to the injection of the thiourea solution. Figure 4.5a compares the XRD
patterns of the nickel sulfide nanocrystals synthesized in the presence and absence of 1-
dodecanethiol under otherwise identical conditions. Powder XRD reveals that the
nanocrystals comprise a mixture of relatively more sulfur-rich cubic Ni3S4 and hexagonal
α-NiS phases in the absence of thiol. In contrast, the injection of either 1.5 or 3.0 mL of 1-
dodecanethiol 5 min after the injection of N,N’-diphenyl thiourea gives phase-pure, sulfur-
deficient, orthorhombic Ni9S8 nanocrystals. Morphologies of the resulting nanocrystals are
comparable in the absence and presence of thiol, suggesting that amount of thiol has a
subtle impact on morphology (Figure 4.5b).

108

Figure 4.5. (a) Powder XRD patterns and (b) TEM micrographs of the products from the
reaction of NiI 2 and 3.0 molar equivalents of N,N’-diphenyl thiourea in the presence and
absence of 1-dodecanethiol (DDT) in oleylamine.

These results agree well with our previous observations on Ni3S2 nanocrystal synthesis
with N,N’-dibutyl thiourea in the presence of 1-dodecanethiol. In both cases, the injection
of 1-dodecanethiol during the growth period of nanocrystals have a substantial effect on
the phase determination. When 1-dodecanethiol is employed along with N,N’-disubstituted
thioureas, it acts as a reactivity-directing agent and favors the formation of sulfur-deficient
Ni3S2 or Ni9S8 phases over relatively more sulfur-rich Ni9S8 or Ni3S4 and α-NiS phases
(Figure 4.1).

Figure 4.6. (a) Powder XRD patterns and (b) TEM micrographs of the products from the
reaction of NiI2 and 3.0 molar equivalents of N,N’-diphenyl thiourea (DPhT) in the
presence of 3.0 mL of 1-dodecanethiol (DDT) in oleylamine.

109

The effect of the injection time of 1-dodecanethiol was also explored (Figure 4.6).
Simultaneous injection of 1-dodecanethiol and N,N’-diphenyl thiourea resulted in the
formation of relatively more polydispersed Ni9S8 nanocrystals with a variety of shapes
including rods, rectangles, quasi-spheres (Figure 4.6). Alternately, larger nanocrystals
with an ill-defined shapes were observed when 1-dodecanethiol was added to the reaction
flask along with oleylamine before the injection of N,N’-diphenyl thiourea (Figure 4.6). In
this case, 1-dodecanethiol decomposed before the injection of N,N’-diphenyl thiourea. By
the time N,N’-diphenyl thiourea is injected at 180
o
C, nucleation and growth of
nanoparticles have already been initiated and 1-dodecanethiol became the primary sulfur
precursor. Larger and polydispersed nanocrystals can be attributed to slower conversion
kinetics in the case of 1-dodecanethiol that leads to a low monomer concentration
subsequently resulting in Ostwald ripening. These control experiments clearly indicate that
injection time of 1-dodecanethiol has no significant impact on the phase of nanocrystals,
i.e., they all crystalize into the same orthorhombic structure. Whereas, the injection time of
1-dodecanethiol primarily dictates morphology of the resulting nanocrystals rather than
their phase.

110

Figure 4.7. XRD patterns of the products from the reaction of NiI2 and 3.0 molar
equivalents of N,N’-diphenyl thiourea at 180 ˚C in the (a) absence and (b) presence of 1-
dodecanethiol under otherwise identical conditions.

To gain insight into the phase evolution of these nanocrystals, aliquots were taken after
15 s, 5 min, 10 min, 20 min, and 60 min of reaction time in the absence and presence of 1-
dodecanethiol, with t = 0 indicating the injection of N,N’-diphenyl thiourea. Thiol was
injected (if applicable) at t = 5 min. Figure 4.7 shows the phase evolution of the nickel
sulfide nanocrystals under these conditions. In the absence of 1-dodecanethiol, the resulting
product is completely amorphous after 15 s. After 5 min, broad diffraction peaks were
observed that might be indexed to the cubic Ni3S4 phase with hexagonal α-NiS impurities
which remained the same until 10 min. After 20 min, a mixture of the Ni 3S4 phase along
with α-NiS phase was still observed with higher intensity, and it remained as a mixed-
phase product through at least 60 min (Figure 4.7a). When an identical reaction was carried
out in the presence of 1-dodecanethiol, a different phase evolution was observed after thiol
addition. The first two aliquots taken at 15 s and 5 min give the same products to the
reaction without thiol, resulting in amorphous products. Immediately after the second
aliquot was taken at 5 min, 1.5 mL of 1-dodecanethiol was injected. In the presence of 1-

111

dodecanethiol, the XRD pattern of the product from the aliquot taken at 10 min can be
indexed to the rhombohedral Ni3S2 phase, and at 20 min a mixture of Ni3S2 and Ni9S8
phases was observed (Figure 4.7b). After 60 min, the product is completely converted into
the orthorhombic Ni9S8 phase, with no evidence of Ni3S2 or any other crystalline Ni–S
phases. This distinction in phase evolution of nanocrystals in the absence and presence of
1-dodecanethiol resembles with the observations of Zhou et al. on Cu2ZnSnS4 (CZTS)
nanocrystals.
35
They have recently reported phase control over CZTS nanocrystals by
employing 1-dodecanethiol along with elemental sulfur and found that an increase in the
amount of 1-dodecanethiol restrains the reactivity of sulfur leading to a phase evolution
from kesterite phase to wurtzite phase. When small amount of 1-dodecanthiol is employed
with elemental sulfur, nanocrystals evolve from Cu1.94S to kesterite CZTS with a
(ZnS)x(Cu2SnS3)1-x intermediate. Whereas, small amount of sulfur along with 1-
dodecanethiol leads to a phase transformation from Cu1.94S to wurtzite phase through a
different (SnS)x(Cu2ZnS3)1-x intermediate.
35
These experiments clearly indicate that 1-
dodecanethiol plays a decisive role in determining both the evolution and resulting phase
of the nickel sulfide nanocrystals and favors the formation of sulfur-deficient Ni9S8 and
Ni3S2 phases over sulfur-rich Ni3S4 and α-NiS phases.

112

Figure 4.8.
1
H NMR of spectra in dichloromethane-d2 of N,N’-diphenyl thiourea heated in
oleylamine at 180 ˚C for 10 min, and then 5 min after 1-dodecanethiol (DDT) was added.
(a) and (b) are showing full spectra and aromatic region only, respectively.
1
H NMR spectra
of starting materials (i.e., oleylamine (OAm) (yellow), N,N’-diphenyl thiourea (DPT)
(pink), and 1-dodecanethiol (DDT) (blue)). Residual solvent is denoted by *.

1
H NMR experiments were carried out to gain further insight into the effect of 1-
dodecanethiol on the reactivity of N,N’-diphenyl thiourea. Identical reaction conditions for
the synthesis of quasi-spherical Ni9S8 nanocrystals were repeated in the absence of NiI2. In
other words, N,N’-diphenyl thiourea was dissolved in oleylamine, and the reaction mixture
was heated to 180 ˚C. The first aliquot was taken 10 min after reaching 180 ˚C, and the
second was taken 5 min after 1-dodecanethiol was injected into the reaction. Figures 4.8
provides the
1
H NMR spectra of aliquots and starting materials (i.e., N,N’-diphenyl
thiourea, oleylamine, and 1-dodecanethiol) in dichloromethane-d2. The
1
H NMR spectra

113

of N,N’-diphenyl thiourea dissolved in oleylamine shows the disappearance of the aromatic
phenyl peaks at 180 ˚C corresponding to the thiourea (  = 7.2-7.5 ppm), indicating thermal
decomposition of N,N’-diphenyl thiourea (Figures 4.8b). Upon addition of 1-
dodecanethiol, there are no apparent changes to the aromatic region of the
1
H NMR
spectrum, further confirming that the thermal decomposition of N,N’-diphenyl thiourea is
independent of 1-dodecanethiol (Figure 4.8b). This suggests that 1-dodecanethiol has no
significant impact on the conversion kinetics of the N,N’-diphenyl thiourea since it is
decomposed and nanocrystal nucleation is initiated before the injection of 1-dodecanethiol.
Formation of sulfur-deficient nickel sulfide phases in the presence of 1-dodecanethiol can
then be attributed to changes in the growth kinetics of the nanocrystals; that is, 1-
dodecanethiol may form complexes with monomers that lead to slower diffusion and
incorporation into growing nanocrystals, and/or it may bind to the nanocrystal surface and
impede monomer incorporation resulting in the formation of more sulfur-deficient
nanocrystals.
28

Figure 4.9. Powder XRD patterns of the products from the reaction of NiI2 and various
amounts of 1-dodecanethiol in oleylamine (OLA).

114

Figure 4.10. TEM micrographs of Ni9S8 nanocrystals synthesized by the reaction of NiI2
with 1-dodecanethiol at 180 ˚C for 2 h with thiol:Ni ratios of (a) 1.5, (b) 3.0 and (c) 6.0.
Histograms of the particle length distributions for nanocrystals synthesized with (d) 1.5,
(e) 3.0, and (f) 6.0 molar equivalents of 1-dodecanethiol.

A new set of reactions were performed to obtain Ni9S8 nanocrystals with different
morphologies; in this case, 1-dodecanethiol was used as the sole sulfur precursor instead
of N,N’-disubstituted thioureas, with the concentration of NiI2 being held constant (Figure
4.1). In a typical reaction, NiI2 was dissolved in 5.0 mL of oleylamine, and varying
amounts of 1-dodecanethiol were injected into the NiI2 solution at 180 ˚C, with the reaction
being allowed to proceed for 2 h (Table 4.2). Figure 4.9 provides the XRD patterns of the
products from the reaction of NiI2 with various amounts of 1-dodecanethiol at 180 ˚C for
2 h. Analysis of the as-synthesized products reveal that the particles crystallize into the
same orthorhombic Ni9S8 structure. The diffraction peaks of the product synthesized using
a molar thiol:Ni ratio of 1.5 are broader than those for ratios of 3.0 and 6.0. A sharpening
of all the diffraction lines was observed upon increasing the thiol:Ni ratio, suggesting an

115

increase in the particle size. Figure 4.10 provides the TEM images of the resulting
nanocrystals synthesized using various amounts of 1-dodecanethiol. These micrographs
indicate that when 1-dodecanethiol is used as the sole sulfur precursor, rod-like Ni9S8
nanocrystals are produced. The nanocrystal lengths were found to be 14.5  3.4, 27.1  6.7
and 62.7  20.6 nm for 1.5, 3.0, and 6.0 molar equivalents of 1-dodecanethiol, respectively
(Table 4.2). The TEM micrographs and particle length distributions show that the length
of the nanorods becomes longer and edge lengths are more polydispersed as the amount of
1-dodecanethiol increased in the reaction.

Table 4.2. Synthetic conditions for the preparation of shape-controlled Ni9S8
nanoparticles. NiI2 is reacted with various sulfur sources in oleylamine at 180 ˚C.

These observations show that 1-dodecanethiol is a key reagent for controlling the phase
and morphology in the binary Ni–S system (Figure 4.1 and Table 4.2). In agreement with
prior literature reports, we found that when 1-dodecanethiol is used as the sole sulfur
precursor, it leads to the formation of orthorhombic Ni9S8 nanocrystals with a rod-like
morphology.
11,23,37
We also found that the nanorod length can be tuned by the amount of
Sulfur precursor S:Ni ratio
Reaction
time (h)
Shape Size (nm)
N,N’-diphenyl thiourea /
1-dodecanethiol
3.0 w/ 3.0 mL of
1-dodecanethiol
4 Spherical 8.8 ± 1.8
N,N’-dibutyl thiourea 10.0 1
Rod-like
aggregates
~ 100
1-dodecanethiol 1.5 2 Rods 14.5 ± 3.4
1-dodecanethiol 3.0 2 Rods 27.1 ± 6.7
1-dodecanethiol 6.0 2 Bricks 62.7 ± 20.6

116

1-dodecanethiol used in the synthesis relative to NiI2, with an increase in the amount of 1-
dodecanethiol resulting in larger particles. This suggests that the ratio of oleylamine to 1-
dodecanethiol is crucial for achieving arrested growth conditions when 1-dodecanethiol is
used as the sole sulfur precursor. When 1-dodecanethiol is employed after the injection of
N,N’-disubstituted thioureas, it acts as a growth-directing agent instead and favors the
formation of sulfur-deficient phases by decreasing the growth kinetics of nanocrystals.
Nearly monodispersed, quasi-spherical Ni9S8 nanocrystals are obtained from the reaction
of NiI2 and N,N’-diphenyl thiourea in the presence of 1-dodecanethiol (Figure 4.1, Figure
4.2 and Table 4.4). When less-reactive N,N’-dibutyl thiourea is used as the sole sulfur
precursor instead of N,N’-diphenyl thiourea, phase-pure Ni9S8 nanocrystals with rod-like
aggregates are obtained in the absence of 1-dodecanethiol (Figure 4.1 and Figure 4.3).
With respect to the different phases obtained by using various N,N’-disubstituted
thioureas, our observations can be attributed to differences in precursor conversion
kinetics. Owen and co-workers demonstrated that the conversion rates of N,N’-
disubstituted thioureas decrease upon the replacement of electron-withdrawing aryl groups
with electron-donating alkyl substituents.
30
Slower conversion kinetics in the case of N,N’-
dibutyl thiourea in the presence of 1-dodecanethiol result in the formation of the most
sulfur-deficient Ni3S2 phase, with the morphology of the final product being dictated by
the amount of 1-dodecanethiol used in the synthesis. In contrast, the relatively more sulfur-
rich Ni3S4 and NiS phases are obtained by using more reactive N,N’-diphenyl thiourea in
the absence of 1-dodecanethiol due to its faster conversion kinetics (Figure 4.1). The
combination of 1-dodecanethiol with the more reactive N,N’-diphenyl thiourea appears to
temper the growth kinetics of nanocrystals and favor formation of the relatively less sulfur-

117

rich Ni9S8 phase (Figure 4.1 and Table 4.2). Tunable conversion kinetics of N,N’-
disubstituted thioureas also enable us to control morphology. Relatively bigger and
polydispersed nanocrystals are obtained when N,N’-dibutyl thiourea is used as the sole
sulfur precursor (Figure 4.2 and 4.3b). Whereas, faster conversion kinetics of N,N’-
diphenyl thiourea result in smaller and nearly monodispersed quasi-spherical nanocrystals
(Figure 4.4). These results are in a good agreement with previous studies. Owen et al. have
reported synthesis of Ni-S nanocrystals from the reaction of nickel (II) stearate with three
different N,N’-disubstituted thioureas (i.e., N-N’-diphenyl thiourea, N-phenyl-N’-n-
dodecyl thiourea, and N-n-hexyl-N’-N’-di-n-butyl thiourea). They found that as-
synthesized nanocrystals become polydispersed and ill-defined as the reactivity of
substituted thioureas decreases. Similarly, they have also observed an increase in the aspect
ratio and volume of CdS nanorods with decreasing reactivities of N,N’-disubstituted
thioureas, stemming from the lower nanocrystal concentration due to the slower conversion
kinetics.
34

4.3.2. Sulfur-rich Ni-S Phases
Effect of temperature and precursor ratio. A series of reactions were carried out to gain
insight into the influence of temperature and precursor ratio on the phase of relatively more
sulfur-rich nickel sulfide nanocrystals by independently varying each parameter and using
more reactive N,N’-diphenyl thiourea as the sole sulfur precursor. First, the possibility of
obtaining different nickel sulfide phases by changing the sulfur to nickel precursor ratio
was investigated. By fixing the reaction time at 4 h and the temperature at 180 ˚C, a series
of reactions were performed with a range of N,N’-diphenyl thiourea:NiI2 molar ratios

118

between 3.0 and 6.0 in the absence of 1-dodecanethiol. The resulting powder XRD patterns
indicate that the sulfur-rich cubic Ni3S4 phase and hexagonal α-NiS are present with a S:Ni
precursor ratio of 3.0 (Figure 4.11a). An increase in the intensities of the reflections at 30,
35, and 46˚ 2θ from α-NiS was observed as the S:Ni precursor ratio increased to 4.0-5.0,
with the intensities of the reflections from the Ni3S4 phase concomitantly decreasing. With
5.0 molar equivalents of N,N’-diphenyl thiourea, α-NiS is favored along with some minor
impurity peaks from Ni3S4. In other words, the cubic Ni3S4 phase appears to convert
hexagonal α-NiS phase by increasing the amount of sulfur precursor. At a S:Ni molar ratio
of 6.0, the resulting nanocrystals exist mainly in the α-NiS phase, along with minor
amounts of Ni3S4 and β-NiS impurities. This set of experiments clearly demonstrates that
the amount of N,N’-diphenyl thiourea plays a significant role in the phase determination of
the nanocrystals, with an increase in the amount of sulfur precursor leading to the
hexagonal α-NiS and rhombohedral β-NiS phases being preferentially formed over cubic
Ni3S4 at 180 ˚C. Simply considering a nickel sulfidation reaction, one would expect the
formation of sulfur-rich Ni3S4 nanocrystals at higher S:Ni ratios. Instead, our results
resemble those of Korgel and Hayakawa, both of whom obtained the sulfur-rich Ni3S4
phase at low sulfur precursor ratios.
22,38
Korgel et al. attributed this result to the faster
sulfidation kinetics stemming from the small size of their nanoparticles.

119

Figure 4.11. Powder XRD patterns of the products from the reaction of NiI2 with variable
amounts of N,N’-diphenyl thiourea: (a) as a function of sulfur S:Ni ratio; all reactions were
carried out at 180 ˚C for 4 h, and (b) as a function of temperature for 4 h with a S:Ni molar
ratio of 4.0.

Next, the role of the temperature on nanocrystal phase determination was explored by
holding the S:Ni molar ratio and reaction time constant. A moderate S:Ni molar ratio of
4.0 was chosen and the reactions were carried out in a range of temperatures between 180
and 250 ˚C for 4 h. The resulting XRD patterns indicate that the nanocrystals synthesized
at 180 and 200 ˚C crystallize into a mixture of the cubic Ni3S4 and hexagonal α-NiS phases
(Figure 4.11b). When the reaction temperature was increased to 220 ˚C, the α-NiS phase
along with some impurities from Ni3S4 and β-NiS were obtained. When the temperature
was further increased to 250 ˚C, the particles mainly crystallize into the rhombohedral β-
NiS phase along with small amounts of α-NiS and Ni3S4 impurities. This set of experiments
demonstrates that temperature also has a significant impact on phase determination; with
increasing reaction temperature, a phase preference from Ni3S4 to α-NiS and then β-NiS
phase was observed. These results are in an agreement with previous reports, where Ni3S4
has previously been shown to be a metastable phase at low temperatures that
disproportionates at higher temperatures to give the NiS and NiS1.03 phases.
39

120

Synthesis of Ni3S4 nanocrystals. Based on these observations regarding the influence of
precursor ratio and reaction temperature, we were able to design synthetic conditions for
the preparation of phase-pure Ni3S4, α-NiS and β-NiS nanocrystals. Ni3S4 is the major
product of the reactions at low S:Ni precursor ratios and low temperatures (Figure 4.11).
Therefore, we reduced S:Ni precursor ratio to 1.5 and characterized the product of the
reaction between NiI2 and 1.5 molar equivalents of N,N’-diphenyl thiourea at 180 ˚C for 4
h (Table 4.1). XRD analysis reveals that the particles crystallize into the cubic polydymite
structure of Ni3S4, with no evidence of α-NiS or β-NiS impurities (Figure 4.12a). The
nanocrystals possess a mean diagonal length of 12.4  2.4 nm measured from ~300
randomly chosen particles (Figure 4.12d). An HR-TEM micrograph of an individual Ni3S4
nanocrystal is given in Figure 4.12c; the nanoparticle appears to be single crystalline with
a lattice spacing of d = 5.5 Å that matches well with the (111) plane of the cubic polydymite
structure.

Figure 4.12. (a) XRD pattern of cubic Ni3S4 nanocrystals. (b) TEM micrograph of 12.4-
nm Ni3S4 nanocrystals. (c) HR-TEM micrograph of an individual Ni3S4 nanocrystal. (d)
Size histogram showing the distribution of nanocrystal diameters for Ni3S4 (N = 301).

121

Synthesis of α-NiS nanocrystals. The first set of experiments on the roles of precursor
ratio and reaction temperature demonstrates that the α-NiS phase is the major product of
the reaction that was performed at 180 ˚C with a S:Ni molar ratio of 5.0 (Figure 4.11).
However, our attempts to obtain phase-pure α-NiS nanocrystals by fine-tuning the
temperature, precursor ratio and reaction time failed. Next, we employed different
coordinating solvents/capping ligands other than oleylamine to prepare phase-pure α-NiS
nanocrystals, inspired by the previous studies showing phase control over Ni–S
nanocrystals by varying the conversion kinetics of precursors with respect to the high-
boiling solvents.
23,25,26
Recently, Revaprasadu et al. reported phase control over colloidal
Ni–S nanoparticles using single-source precursors in primary amine solvents.
26
In their
study, dodecylamine, hexadecylamine and oleylamine led to the formation of Ni3S4, Ni3S2,
and mixed phase particles, respectively, under otherwise identical conditions, which may
be attributed to the difference in the precursor conversion kinetics of the single-source
precursor as a function of amine solvent. Similarly, Owen et al. reported a difference in the
conversion kinetics for N-phenyl-N’-dodecyl thiourea in different high-boiling solvents.
The conversion rate constant of N-phenyl-N’-dodecyl thiourea was shown as higher in tri-
n-butyl amine when compared to that of in oleic acid.
34

In order to prepare phase-pure α-NiS, we kept S:Ni ratio of 5.0 constant and used
dodecylamine in place of oleylamine as the coordinating solvent. N,N’-diphenyl thiourea
was quickly injected into the NiI2 solution at 180 ˚C, and the reaction was maintained at
that temperature for 5 min (Table 4.1). Powder XRD analysis reveals that the product is
composed of hexagonal α-NiS with no evidence of any Ni3S4 or β-NiS impurities (Figure
4.13a). The resulting nanocrystals are quasi-spherical, with an average particle diameter of

122

8.9  2.4 nm determined by TEM analysis (Figure 4.13b-d). An HR-TEM image of an
individual nanocrystal is shown in Figure 4.13c, which suggests that the particle is single-
crystalline with a measured interplanar spacing of d = 2.5 Å being in good agreement with
the d-spacing of the (101) planes of hexagonal α-NiS. According to bulk phase diagram of
nickel sulfide, the α-NiS phase is the high-temperature polymorph with the β-NiS phase
being thermodynamically preferred;
1,19
however, on the nanoscale, a kinetically favored
metastable α-NiS phase can be formed at lower temperatures.
24

Figure 4.13. (a) XRD pattern of hexagonal α-NiS nanocrystals. (b) TEM micrograph of
8.9-nm α-NiS nanocrystals. (c) HR-TEM micrograph of an individual α-NiS nanocrystal.
(d) Size histogram showing the distribution of nanocrystal diameters for α-NiS (N = 300).

Synthesis of β-NiS nanocrystals. We further attempted to obtain phase-pure β-NiS
nanocrystals. In the series of reactions as a function of S:Ni ratio and temperature, β-NiS
impurities were observed at a S:Ni ratio of 6.0 and the reaction temperature of 180 ˚C.
Similarly, when the reaction temperature was increased to 220 ˚C, the β-NiS phase was
observed to be the major product with a S:Ni ratio of 4.0 (Figure 4.11). Therefore, we tried
to fully convert α-NiS nanocrystals to the thermodynamically favored β-NiS phase by the

123

aid of increasing reaction temperature and S:Ni precursor ratio. The reaction of NiI2 with
an 8.0 molar excess of N,N’-diphenyl thiourea in oleylamine at 220 ˚C for 4 h gives a
product primarily indexed to the millerite β-NiS phase with a rhombohedral structure by
XRD (Table 4.1); however, some minor impurity peaks from hexagonal α-NiS were still
found (Figure 4.14a).
The reaction of NiI2 with an 8.0 molar excess of N,N’-phenyl thiourea in oleylamine at
220 ˚C for 4 h gives a colloidally unstable product with large particles and XRD analysis
of the product reveals β-NiS nanocrystals with minor α-NiS impurities (Figure 4.14b).
Although hexagonal α-NiS nanocrystals can be converted into rhombohedral β-NiS phase
under forcing conditions, we exceeded kinetic control regime and lost control over
morphology under these conditions. In contrast to other Ni-S phases synthesized by kinetic-
control, β-NiS phase was obtained under thermodynamic control.

Figure 4.14. (a) Powder XRD patterns of rhombohedral β-NiS nanocrystals synthesized
by N,N’-diphenyl thiourea with the α-NiS impurity shown by (*). (b) TEM micrograph of
β-NiS nanocrystals.

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4.4. Experimental
4.4.1. General Considerations
Nickel (II) iodide (NiI2, Alfa Aesar, 99.5%), N,N’-diphenyl thiourea (Alfa Aesar, 98%),
N,N’-dibutyl thiourea (Alfa Aesar, 98%), oleylamine (cis-9-octadecenylamine, Sigma
Aldrich, 70%), dodecylamine (1-aminododecane, Sigma Aldrich, 98%), dibenzylamine
(Alfa Aesar, 98%) and 1-dodecanethiol (Sigma Aldrich, 98%) were all purchased and used
without further purification. Nanoparticle syntheses were conducted under N2 using
Schleck techniques in the absence of water and oxygen.
4.4.2. Ni3S2 Nanocrystal Synthesis
NiI2 (0.19 mmol, 0.06 g) and degassed oleylamine (15.2 mmol, 5.0 mL) were added to
a three-neck flask fitted with a reflux condenser and rubber septa. The solution was heated
to 120 ˚C and degassed for 30 min under vacuum. N,N’-dibutyl thiourea (1.9 mmol, 0.36
g) was dissolved in dibenzylamine (15.6 mmol, 3.0 mL) and the solution was sparged by
bubbling nitrogen through it for 15 min. The solution of NiI2 in oleylamine was heated to
180 ˚C and then the solution of N,N’-dibutyl thiourea was quickly injected into the flask
under flowing N2. After 5 min, 1-dodecanethiol (12.5 mmol, 3.0 mL) was also injected into
the reaction mixture and the reaction was allowed to proceed for 1 h with stirring under
flowing N2. The reaction was quenched by placing it in a water bath and allowing it to cool
to room temperature.
4.4.3. Ni9S8 Nanocrystal Synthesis
In a typical synthesis, NiI2 (0.38 mmol, 0.12 g) and degassed oleylamine (15.2 mmol,
5.0 mL) were added to a three-neck flask fitted with a reflux condenser and rubber septa.
The solution was heated to 120 ˚C and degassed for 30 min under vacuum to eliminate

125

adventitious water and dissolved oxygen. N,N’-diphenyl thiourea (1.14 mmol, 0.26 g) was
dissolved in dibenzylamine (10.4 mmol, 2.0 mL) and the solution was sparged by bubbling
N2 through it for 15 min. The solution of NiI2 in oleylamine was heated to 180˚C under
flowing N2 and then the N,N’-diphenyl thiourea solution in dibenzylamine was quickly
injected into the reaction flask. After 5 min, 1-dodecanethiol (12.5 mmol, 3.0 mL) was
subsequently injected into the reaction mixture and then the reaction was allowed to
proceed at 180 ˚C for 4 h with stirring under flowing N2. After 4 h, the reaction was
thermally quenched by placing it in a water bath and allowing it to cool to room
temperature.
Ni9S8 nanorods were synthesized by a similar synthetic approach. In this case, 1-
dodecanethiol was used as the sole sulfur precursor. In a typical reaction, 0.38 mmol (0.12
g) of NiI2 was dissolved in 5.0 mL of oleylamine. Various amounts of 1-dodecanethiol
(i.e., 1.5, 3.0 and 6.0 molar equivalents) were rapidly injected into the solution of NiI2 in
oleylamine at 180 ˚C. The reaction mixture was allowed to react for 2 h with stirring under
flowing N2, followed by thermally quenching the reaction by placing it in a water bath and
allowing it to cool to room temperature.
4.4.4. Ni3S4 Nanocrystal Synthesis
NiI2 (0.38 mmol, 0.12 g) and degassed oleylamine (15.2 mmol, 5.0 mL) were added to
a three-neck flask fitted with a reflux condenser and rubber septa. The solution was heated
to 120 ˚C and degassed for 30 min under vacuum. N,N’-diphenyl thiourea (0.57 mmol, 0.13
g) was dissolved in dibenzylamine (7.8 mmol, 1.5 mL) and the solution was sparged by
bubbling N2 through it for 15 min. The solution of NiI2 in oleylamine was heated to 180
˚C, and then the N,N’-diphenyl thiourea solution in dibenzylamine was quickly injected

126

into the reaction flask and allowed to react for 4 h with stirring under flowing N 2. The
reaction was quenched by placing it in a water bath and allowing it to cool to room
temperature.
4.4.5. α-NiS Nanocrystal Synthesis
NiI2 (0.38 mmol, 0.12 g) and degassed dodecylamine (21.7 mmol, 5.0 mL) were added
to a three-neck flask fitted with a reflux condenser and rubber septa. The solution was
cycled between vacuum and N2 several times at room temperature. N,N’-diphenyl thiourea
(1.9 mmol, 0.43 g) was dissolved in dibenzylamine (15.6 mmol, 3.0 mL) and the solution
was sparged by bubbling N2 through it for 15 min. The solution of NiI2 in dodecylamine
was heated to 180˚C, and then the N,N’-diphenyl thiourea solution was quickly injected
into the reaction flask and allowed to react for 5 min with stirring under flowing N 2. The
reaction was quenched by placing it in a water bath and allowing it to cool to room
temperature.
4.4.6. β-NiS Nanocrystal Synthesis
NiI2 (0.19 mmol, 0.06 g) and degassed oleylamine (15.2 mmol, 5.0 mL) were added to
a three-neck flask fitted with a reflux condenser and rubber septa. The solution was heated
to 120 ˚C and degassed for 30 min under vacuum. N,N’-dibutyl thiourea (3.8 mmol, 0.72
g) was dissolved in dibenzylamine (20.8 mmol, 4.0 mL) and the solution was sparged by
bubbling N2 through it for 15 min. The solution of NiI2 in oleylamine was heated for 220
˚C, and then the N,N’-dibutyl thiourea solution was quickly injected into the reaction flask
and allowed to react for 4 h with stirring under flowing N2. The reaction was quenched by
placing it in a water bath and allowing it to cool to room temperature.

127

4.4.7. Nanocrystal Purification
Nanocrystals were purified by precipitation in 25 mL of ethanol followed by
centrifugation at 6000 rpm for 10 min. The supernatant was discarded. The precipitated
nanocrystals were redispersed in 20 mL of hexanes and centrifuged again at 6000 rpm for
3 min, causing the larger particulates to settle. The precipitate was discarded, and the
nanocrystals suspended in hexanes were reprecipitated again by addition of 20 mL of
ethanol. The precipitated nanocrystals were finally dispersed in either hexanes or toluene.
All workup procedures were carried out in air.
4.4.8. Instrumentation
Powder X-ray diffraction (XRD) analyses were performed on Rigaku Ultima IV X-ray
diffractometer operated at 44 mA and 40 kV using a Cu Kα radiation source (λ = 1.5406
Å). Diffraction patterns were collected in the 2θ range of 10˚ to 70˚. The step size and
collection time were 0.008˚ and 1 s per step, respectively. All patterns were recorded under
ambient conditions. Phases were assigned by the powder diffraction files of (PDFs) of the
International Center Diffraction Data (ICDD) using Jade 9.0 software.
Transmission electron microscopy (TEM) imaging was performed on a JEOL JEM-
2100 microscope at an operating voltage of 200 kV, equipped with a Gatan Orius CCD
camera. Samples for TEM studies were prepared by drop-casting a stable suspension of
nanocrystals in toluene on a 200 mesh Cu grid coated with a lacey carbon film (Ted Pella,
Inc.).
Nuclear magnetic resonance (NMR) spectra were collected on a 600 MHz Varian
spectrometer, and the data was analyzed using MestReNova version 12.0.0 software.
Samples were prepared in either chloroform-d or dichloromethane-d2.

128

4.5. Conclusions
An analysis of key reaction parameters enabled a high degree of control over the phase,
composition, and morphology of colloidal nickel sulfide nanocrystals. A series of well-
defined Ni3S4, α-NiS, Ni9S8, and Ni3S2 nanocrystals were prepared via solution-phase
syntheses using N,N’-disubstituted thioureas as the sulfur precursor. In addition to choice
of the primary sulfur precursor, a combination of additional experimental parameters (i.e.,
solvent and ligand type, precursor ratio, and reaction temperature and time) were needed
to exert phase and morphology control. When 1-dodecanethiol is used as a reactivity-
directing agent, a series of sulfur-deficient Ni9S8 and Ni3S2 phases with spherical, rod-like,
and multipod-type morphologies are attained. Both the reaction temperature and S:Ni
precursor ratio were demonstrated as being key parameters for the phase evolution of the
more sulfur-rich Ni3S4, α-NiS, β-NiS phases when N,N’-diphenyl thiourea was used as the
sole sulfur precursor in the absence of 1-dodecanethiol. At a S:Ni ratio of 1.5 and a reaction
temperature of 180 ˚C, phase-pure Ni3S4 nanocrystals were obtained by using oleylamine
as the solvent and capping ligand. In order to obtain phase-pure α-NiS nanocrystals,
dodecylamine was used as coordinating solvent instead of oleylamine, and the S:Ni
precursor ratio was increased to 5.0. A further increase in the S:Ni ratio and reaction
temperature results in the transformation of α-NiS phase into thermodynamically favorable
β-NiS phase. The strategy used to prepare phase-pure nickel sulfide nanocrystals with well-
defined morphologies may facilitate the development of this significant class of nano-
structured materials for energy-related applications and catalysis.

129

4.6. References

(1) Kullerud, G.; Yund, R. A. The Ni-S System and Related Minerals. J. Petrol. 1962,
3, 126–175.

(2) Manthiram, A.; Jeong, Y. U. Ambient Temperature Synthesis of Spinel Ni3S4: an
Itinerant Electron Ferrimagnet. J. Solid State Chem. 1999, 147, 679-681.

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Chapter 5. Synthesis and Characterization of Thiospinel CoNi2S4 and FeNi2S4
Nanocrystals*

*Unpublished results

5.1. Introduction

Transition metal sulfides that crystallize into a cubic spinel structure are known as
thiospinels. The spinel structure can be described by the general formula of AB2X4, where
A and B are typically divalent and trivalent metal cations, respectively, and X can be either
oxygen or sulfur. In the normal spinel structure, divalent A metal cations and trivalent B
metal cations occupy tetrahedral and octahedral positions, respectively, whereas, anions sit
at the polyhedral vertexes.
1,2
Since anions usually adopt the -2 oxidation state, the A and B
cations can also adopt +4 and +2 oxidations states, respectively, and in this case, the
structure is called as inverse spinel. Various distributions of A and B metal cations also
exist with different ratios occupying both tetrahedral and octahedral holes, which is known
as complex spinel. Cation radii, Coulomb interactions between cations, and crystal field
effects of the octahedral site preference energy are the main factors that determine the
distribution of metal cations in the spinel structure.
3

Single-metal thiospinels, such as Ni3S4, Co3S4 and Fe3S4, have been widely
investigated for their electronic, magnetic, and catalytic properties. These end members
can also form solid-solutions. In the Ni3S4-CoNi2S4 system, some portion of nickel is
substituted by cobalt to form CoNi2S4. CoNi2S4 (siegenite) is a normal spinel with divalent
cobalt and trivalent nickel cations occupying tetrahedral and octahedral sites, respectively.
Similarly, FeNi2S4 (violarite) can be prepared by the substitution of nickel by iron in the
Ni3S4-FeNi2S4 system. Both CoNi2S4 and FeNi2S4 display interesting magnetic and

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electrochemical properties and have been studied for various electrochemical applications.
including electrocatalysts for hydrogen evolution and oxygen evolution reactions,
electrodes for rechargeable ion batteries, and supercapacitors.
5–13

Conventionally, bulk thiospinels are synthesized by the solid-state reaction of metals
with sulfur in an evacuated tube
14
or the reaction of oxides with H2S
15
at temperatures
above 500 ˚C. Recently, syntheses of CoNi2S4 and FeNi2S4 thiospinels at lower
temperatures through solvothermal
8,16–18
, hydrothermal
19
, fast solidification
20
and wet-
chemical approaches
21
have also been reported. Although these synthetic methods are more
viable to obtain high surface area materials for energy applications, the resulting particles
are usually in micro-scale in dimension with ill-defined morphologies. In 2016, the Schaak
group demonstrated low temperature solution-phase synthesis of thiospinel CuCo2S4
nanocrystals by a heating up method.
22
10-nm colloidal CuCo2S4 nanocrystals were
obtained by the reaction of Ni(acac)2 and Co(acac)2 with elemental sulfur in oleylamine at
200 ˚C in the presence of oleic acid and tri-n-octylphosphine oxide. Colloidal CuCo2S4
nanocrystals were shown to be highly active materials toward oxygen evolution reaction
under alkaline conditions with an overpotential of 395 mV to produce a current density of
10 mA cm
−2
.
Herein, we demonstrate the low temperature solution-phase synthesis of colloidal
CoNi2S4 and FeNi2S4 nanocrystals with highly well-defined morphologies by a hot-
injection approach. Ni(acac)2, Co(acac)2 and 1-dodecanethiol are used as precursors and
oleylamine is used as high boiling solvent and capping ligand in the synthesis. Structural
and morphological characterization of as-synthesized nanocrystals by powder X-ray
diffraction and transmission electron microscopy will also be presented.

135

5.2. Experimental Details

5.2.1. General Considerations
Nickel (II) iodide (NiI2, Alfa Aesar, 99.5%), nickel(II) acetylacetonate (Ni(acac)2, Alfa
Aesar, 95%, anhydrous), cobalt(II) acetylacetonate (Ni(acac)2, Sigma Aldrich, 97%,
anhydrous), iron (II) iodide (FeI2, Strem Chemicals, 99.5%), N,N’-diphenyl thiourea (Alfa
Aesar, 98%), N,N’-dibutyl thiourea (Alfa Aesar, 98%), oleylamine (cis-9-
octadecenylamine, Sigma Aldrich, 70%), dibenzylamine (Alfa Aesar, 98%) and 1-
dodecanethiol (Sigma Aldrich, 98%) were all purchased and used without further
purification. Nanoparticle syntheses were conducted under N2 using Schleck techniques in
the absence of water and oxygen.

5.2.2. Nanocrystal Syntheses

5.2.2.1. Synthesis of CoNi2S4 Nanocrystals

Ni(acac)2 (0.38 mmol, 0.097 g), Co(acac)2 (0.19 mmol, 0.05 g), and degassed
oleylamine (15.2 mmol, 5.0 mL) were added to a three-neck flask fitted with a reflux
condenser and rubber septa. The solution was heated to 120 ˚C and degassed for 30 min
under vacuum. The solution of Ni(acac)2 and Co(acac)2 in oleylamine was heated to 180
˚C and then the solution of 1-dodecanethiol (1.04 mmol, 0.25 mL) was quickly injected
into the flask under flowing N2. The reaction was allowed to proceed for 1 h with stirring
under flowing N2. The reaction was quenched by placing it in a room-temperature water
bath and allowing it to cool. After being cooled to room temperature, the reaction mixture
was precipitated with 30 mL of ethanol, sonicated, and centrifuged (6000 rpm for 10 min)
to yield a black solid. The black solid was dispersed in 5 mL of hexanes and then
reprecipitated with 30 mL of ethanol. Dispersion/precipitation was repeated three times
with hexanes (5 mL) and ethanol (30 mL) to yield the purified product.

136

5.2.2.2. Synthesis of FeNi2S4 Nanocrystals

NiI2 (0.38 mmol, 0.12 g), FeI2 (0.19 mmol, 0.06 g), and degassed oleylamine (15.2
mmol, 5.0 mL) were added to a three-neck flask fitted with a reflux condenser and rubber
septa. The solution was heated to 120 ˚C and degassed for 30 min under vacuum. The
solution of NiI2 and FeI2 in oleylamine was heated to 180 ˚C and then the solution of 1-
dodecanethiol (1.04 mmol, 0.25 mL) was quickly injected into the flask under flowing N2.
The reaction was allowed to proceed for 2 h with stirring under flowing N2. The reaction
was quenched by placing it in a room-temperature water bath and allowing it to cool. After
being cooled to room temperature, the reaction mixture was precipitated with 30 mL of
ethanol, sonicated, and centrifuged (6000 rpm for 10 min) to yield a black solid. The black
solid was dispersed in 5 mL of toluene and then reprecipitated with 30 mL of ethanol.
Dispersion/precipitation was repeated three times with toluene (5 mL) and ethanol (30 mL)
to yield the purified product.

5.2.3. Characterization of Nanocrystals
Powder X-ray diffraction (XRD) analyses were performed on a Rigaku Ultima IV X-
ray diffractometer using a Cu Kα radiation source (λ = 1.54 Å). Transmission electron
microscopy (TEM) was performed on a JEOL JEM-2100 microscope at an operating
voltage of 200 kV, equipped with a Gatan Orius CCD camera.

137

5.3. Results and Discussion

5.3.1. CoNi2S4 Nanocrystals

Highly crystalline CoNi2S4 nanocrystals were prepared via the fast injection of 1-
dodecanethiol into a solution of Co(acac)2 and Ni(acac)2 in oleylamine at 180 ˚C.
Immediately after the injection of 1-dodecanethiol, black particles were observed
indicating fast nucleation and the mixture was allowed to react for 1 h. The reaction mixture
was rapidly cooled down to room temperature with a water bath. After cooling down to
room temperature, the reaction mixture was precipitated by 30 mL of ethanol and
centrifuged at 6000 rpm for 10 min to yield a black solid. Black solid was then dispersed
in 5 ml of hexanes and reprecipitated by 30 mL of ethanol to yield a purified product. The
powder X-ray diffraction (XRD) patterns of the purified product is shown in Figure 5.1.
Diffraction lines can be indexed to cubic spinel structure. TEM micrographs of the as-
synthesized nanocrystals display a highly uniform, nearly monodispersed spherical
morphology.

Figure 5.1. (a) Powder XRD patterns of cubic spinel CoNi2S4 nanocrystals synthesized by
1-dodecanethiol. (b) TEM micrograph of CoNi2S4 nanocrystals.

138

To have a better insight into the evolution of nanocrystals, aliquots were taken 5, 10,
20, 30, 50, and 70 min after the injection of 1-dodecanethiol. Figure 5.2 indicates phase
evolution of CoNi2S4 nanocrystals. The first aliquot taken 5 min after the injection of 1-
dodecanethiol is completely amorphous with highly intense low angle peaks indicating
presence of polysulfides in the reaction mixture. At 5 min, the reddish black color of the
reaction mixture also indicates large amount of unreactive polysulfides, suggesting
decomposition and slow conversion of 1-dodecanethiol. At 10 min and 20 min, the
intensities of low angle peaks decrease indicating decomposition of polysulfides, but the
resulting nanoparticles are still amorphous. After 30 min, low angle polysulfide peaks
completely disappear, and nanoparticles start to crystallize into a cubic spinel structure.

Figure 5.2. XRD patterns of the aliquots from the reaction of Co(acac)2 and Ni(acac)2 in
oleylamine and 1-dodecanethiol at 180 ˚C at various reaction times.

139

Cubic spinel CoNi2S4 nanoparticles can also be synthesized by N,N’-diphenyl thiourea
as the sulfur precursor. The same amounts of the metal precursors as in the previous case
were dissolved in oleylamine, and four molar equivalence of N,N’-diphenyl thiourea with
respect to cobalt precursor was injected at 180 ˚C. Aliquots were taken at 5, 10, 20, 30, 40,
and 60 min after the injection of N,N’-diphenyl thiourea. Figure 5.3 shows powder X-ray
diffraction patterns of CoNi2S4 nanocrystals taken at different reaction times under
otherwise identical conditions. Aliquots taken at 5, 10, and 20 min indicate poorly
crystalline products, but after 40 min highly crystalline cubic spinel CoNi 2S4 nanocrystals
can be obtained with various morphologies (Figure 5.3b).

Figure 5.3. (a) XRD patterns of the aliquots from the reaction of Co(acac)2, Ni(acac)2 and
N,N’-diphenyl thiourea in oleylamine at 180 ˚C with various reaction times. (b) TEM
micrographs of nanoparticles taken from different aliquots.

140

5.3.2. FeNi2S4 Nanocrystals

FeNi2S4 nanocrystals were synthesized by a similar low temperature solution-phase
approach. In a typical synthesis, iron(II) iodide (0.19 mmol, 0.06 g) and nickel(II) iodide
(0.38 mmol, 0.12 g) were dissolved in 5 mL of oleylamine at 180 ˚C. 0.25 mL of 1-
dodecanethiol was rapidly injected at 180 ˚C and allowed to react for 2 h. The reaction
mixture was quickly quenched with a water bath and a greenish black slurry was observed.
30 mL of ethanol was added to the reaction mixture, sonicated for 15 min and centrifuged
at 6000 rpm for 10 min to precipitate particles. Sticky black particles were then dispersed
in 5 mL of toluene and reprecipitated by 30 mL of ethanol to obtained purified nanocrystals.
Figure 5.4 shows powder X-ray diffraction patterns of as-purified nanocrystals. Diffraction
patterns can be indexed to a cubic spinel structure without any iron oxide or metal sulfide
impurities.

Figure 5.4. (a) Powder XRD patterns of cubic spinel FeNi2S4 nanocrystals synthesized by
1-dodecanethiol. (b) TEM micrograph of FeNi2S4 nanocrystals.

141

FeNi2S4 nanocrystals can also be synthesized by using N,N’-diphenyl thiourea as sulfur
precursor. In a typical synthesis, iron(II) iodide (0.19 mmol, 0.06 g) and nickel(II) iodide
(0.38 mmol, 0.12 g) were dissolved in oleylamine (15.2 mmol, 5.0 mL) at 180 ˚C. N,N’-
diphenyl thiourea (0.8 mmol, 0.183 g) was dissolved in dibenzylamine (10.4 mmol, 2.0
mL) the solution was sparged by bubbling N2 through it for 15 min. N,N’-diphenyl thiourea
solution was then quickly injected into reaction flask at 180 ˚C and allowed to react at least
1 h. After rapid cooling down to room temperature by a water bath, 30 mL of ethanol was
added to the reaction mixture to precipitate particles. A black solid was obtained by
centrifuging at 6000 rpm for 10 min, and redispersed in 5 mL of hexanes.
Precipitation/dispersion cycle was repeated three times for purification.

Figure 5.5. (a) Powder XRD patterns of cubic spinel FeNi2S4 nanocrystals synthesized by
N,N’-diphenyl thiourea with various reaction times. (b) TEM micrograph of FeNi2S4
nanocrystals with 3 h reaction.

142

Figure 5.5 displays powder XRD patterns as-synthesized nanocrystals by N,N’-
diphenyl thiourea with various reaction times. XRD patterns can be indexed cubic spinel
structure, but an additional peak at 35° was also found indicating the presence of iron oxide
impurities in the product. As a control, XRD patterns of a 4 h reaction were taken after
leaving the diffraction plate under ambient conditions overnight. An increase in the number
of impurity peaks as well as in their intensities was observed. This might be attributed
oxidation of iron on the nanocrystal surface rather than formation of iron oxide during the
reaction. In the case of 1-dodecanethiol, the absence of iron oxide impurities might be
related to polysulfides and other organic species that cover the surface of nanocrystals and
potentially blocks the exposure of surface irons to the air.

5.4. Conclusions

Mixed-metal thiospinel CoNi2S4 and FeNi2S4 nanocrystals were synthesized by a low
temperature solution-based synthetic approach. Metal acetylacetonates in oleylamine
reacted with 1-dodecanethiol at 180 ˚C for 1 h to yield well-defined, highly crystalline
cubic spinel CoNi2S4 nanocrystals. In the case of FeNi2S4 nanocrystal synthesis, iron(II)
iodide and nickel(II) iodide in oleylamine along with 1-dodecanethiol were used as
precursors. When N,N’-diphenyl thiourea was used in place of 1-dodecanethiol as sulfur
precursor iron oxide impurities were also found in the product. These solution-based
synthetic approaches may be extended to the preparation of other thiospinel nanocrystals.

143

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Abstract (if available)

Abstract Electrochemical, magnetic, and electronic properties of functional nanomaterials depend on their size, shape, composition, and crystal phase. Therefore, having a high degree of control over the synthesis of functional nanomaterials is significant for practical applications. The majority of synthetic approaches for functional nanomaterials require highly energy- and capital-intensive conditions and, as such, there is a need to develop new methodologies for the synthesis of functional nanomaterials under more benign conditions. In this work, large-scale, solution-based syntheses of anti-NASICON Fe₂(MoO₄)₃ and Fe₂(WO₄)₃ nanoparticulates, as well as synthetic control over nanocrystal phase in the complicated binary nickel–sulfur phase space, have been featured. ❧ The correlation of structural changes in the frameworks of anti-NASICON type polyanionic cathode materials with their electrochemical properties has been demonstrated through a wide spectrum of structural and electrochemical characterization techniques. Two-phase structural change upon lithium insertion into Fe₂(MO₄)₃, where M = Mo or W, has been confirmed by in situ powder X-ray experiments and galvanostatic electrochemical cycling. Whereas, sodium insertion into Fe₂(MoO₄)₃ proceeds trough a single-phase, solid-solution type mechanism. To gain better insight on the mechanistic difference between lithium and sodium insertion into Fe₂(MoO₄)₃ a combination of high-resolution diffraction techniques and neutron total scattering techniques have been used. Symmetry-mode analysis demonstrated that concerted polyhedral rotations facilitates lithium and sodium insertion into these polyanionic cathode materials. In the case of lithium insertion, the monoclinic pristine structure transforms into a lithiated orthorhombic structure via the polyhedral rotational distortions due the electrostatic interactions between oxygens in the polyhedral subunits and inserted lithium ions. Although sodium insertion into Fe₂(MoO₄)₃ is fundamentally similar, the larger ionic radius of sodium restricts concerted polyhedral rotations and results in a less ordered sodiated structure. ❧ Independent preparation of morphologically well-defined nanocrystals of Ni₃S₄, NiS, Ni₉S₈, and Ni₃S₂ has also been presented via a systematic evaluation of factors that simultaneously control the phase and yield well-defined nickel sulfide nanocrystals. Conversion kinetics of N,N’-disubstituted thioureas have been found as the main factor for the nanocrystal phase control in the binary Ni-S system. Relatively less reactive N,N’-butyl thiourea leads to formation of the sulfur-deficient Ni₉S₈ phase, whereas, more reactive N,N’-phenyl thiourea favors stoichiometric NiS and sulfur-rich Ni₃S₄ phases. Employment of 1-dodecanethiol during the nanocrystal growth stage impede sulfur incorporation into growing nanocrystals and most sulfur-deficient Ni₃S₂ has been obtained by using N,N’-butyl thiourea in the presence of 1-dodecanethiol. In the absence of 1-dodecanethiol, a phase evolution from Ni₃S₄ to NiS phase with an increase in the temperature and sulfur to nickel precursor ratio by using N,N’-phenyl thiourea as the primary sulfur precursor was confirmed. ❧ This synthetic approach has then been expanded to mixed-metal thiospinels. Cubic spinel CoNi₂S₄ and FeNi₂S₄ nanoparticles have been synthesized by the fast-injection of 1-dodecanethiol into a hot mixture of metal precursors in oleylamine at 180℃. Structural and morphological characterization of as-synthesized nanocrystals has been demonstrated as well.

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Creator Barim, Gözde (author)

Core Title Solution phase synthesis routes to functional nanomaterials for energy storage

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 11/07/2018

Defense Date 10/15/2018

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

Tag chalcogenides,energy storage,intercalation,nanocrystals,NASICON,OAI-PMH Harvest,phase control,rechargeable ion batteries,solution-phase synthesis,thiospinels

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Advisor Brutchey, Richard L. (committee chair), Feakins, Sarah J. (committee member), Marinescu, Smaranda C. (committee member)

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