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Ph.D. Dissertation

A Dual-Stage Hydrothermal Flow Reactor

for Green and Sustainable Synthesis of

Advanced Hybrid Nanomaterials

Aarhus University

Department of Chemistry and iNANO

Henrik L. Hellstern

2016

i

Preface

This dissertation has been submitted to the Faculty of Science and

Technology as partial fulfillment of the requirements to obtain the Ph.D.

degree in Nanoscience. All research was conducted under the supervision of

Prof. Bo Brummerstedt Iversen (main supervisor) and Dr. Martin Bremholm

(co-supervisor) at Center for Materials Crystallography at Department of

Chemistry and at the Interdisciplinary Nanoscience Center, iNANO from

November 2012 to February 2016.

The dissertation describes the construction of a dual-stage hydrothermal flow

reactor, and the nanocomposites synthesized in it. To preserve this main

theme several side-projects have been omitted including construction of a

microfluidic droplet reactor, work on the PtGd system, and syntheses of

platinum supported on C3N4 platelets.

For layout purposes certain pages have been left as blank

ii

Acknowledgements

First and foremost I would like to thank my supervisors Professor Bo

Brummerstedt Iversen and Assistant Professor Martin Bremholm for

entrusting me with this project and for their expert guidance. I am deeply

grateful for being given the opportunity to work with the exciting and

challenging field of hydrothermal flow synthesis and to form a part of the

CMC group.

Peter Hald and Jacob Becker are gratefully acknowledged for their help with

reactor design and construction. Drawing from their experiences with flow

reactors has been central to my reactor work.

I owe special thanks to Ander Mørup and Per Christensen for our many

discussions on high-temperature/high-pressure flow reactor setups and

LabVIEW programming. They are furthermore thanked for providing a fun

atmosphere in the office throughout the years including its changing crew of

Yanbin Shen, Maika Klemmer, Martin Søndergaard, Jacob Ahlburg and Niels

Broge.

Everyone, who used the reactor as part of their projects are acknowledged for

their collaboration including Martin Bondesgaard, Sanna Sommer, Mathilde

Múzquiz, and Pedro Magalhães.

Aref Mamakhel is owed special thanks for extensive discussions on

nanoparticle synthesis and for assistance with electron microscopy. It is a

pleasure to work with someone with such a vast synthesis experience.

Mogens Christensen is thanked for many discussions on X-ray diffraction and

on hydrothermal synthesis of magnetic nanoparticles.

Britta Lundtoft and Bo Richter are acknowledged for maintaining the

scientific instruments and lab equipment which have been paramount to my

investigations.

iii

Reactor construction holds many practical challenges, and Erik Ejler from the

chemical workshop is owed special thanks for all his help. It is indispensable

to be able to receive swift help with day-to-day problems.

Tommy Kessler is thanked for assistance with the electrical work. Somehow,

when things didn’t work, he fixed it, and I wasn’t always entirely sure how.

Jacques Chevallier is thanked for assistance with electron microscopy.

Everyone at the section of Inorganic Chemistry is thanked for providing a

great research environment as well as a fun working space.

Camilla Hellstern and Maika Klemmer are gratefully acknowledged for proof-

reading.

I would furthermore like to thank the Faculty of Science and Technology,

CMC and Innovation Fund Denmark for funding my studies.

Finally, I would like to thank Camilla. For everything.

Henrik Lund Hellstern

Aarhus University

January 2016

iv

List of Publications

I. Hellstern, H., Becker, J., Hald, P., Bremholm, M., Mamakhel, A.,

Iversen, B., Development of a Dual-Stage Continuous Flow Reactor for

Hydrothermal Synthesis of Hybrid Nanoparticles, Ind. Eng. Chem.

Res., 2015, 54 (34), 8500-8508

II. Hellstern, H., Bremholm, M., Mamakhel, A., Becker, J., Iversen, B.,

Hydrothermal Synthesis of TiO2@SnO2 Hybrid Nanoparticles in a

Continuous Flow Dual-Stage Reactor, ChemSusChem, 2016 (accepted)

III. Hellstern, H., Bremholm, M., Mamakhel, A., Iversen, B., Core-shell

Nanoparticles by Silica Coating of Metal Oxides in a Dual-Stage

Hydrothermal Flow Reactor, Chem. Commun., 2016 (accepted)

IV. Bondesgaard, M., Becker, J., Xavier, J., Hellstern, H., Mamakhel, A.,

Iversen, B., Guide to by-products formed in organic solvents under

solvothermal conditions, J. Supercrit. Fluids (submitted)

V. Hellstern, H., Iversen, B., Abkey, Ü., Surface Characterization of TiO2

Nanoparticles via Hyperpolarized Solid-State NMR (in preparation)

v

Abstract

Nanocomposites are a group of materials of growing scientific interest. The

combination of two different materials into a single hybrid particle on the

nanoscale can result in multifunctional materials or be used to enhance

existing properties through synergistic interactions. Such novel materials can

be synthesized hydrothermally in a dual-stage flow reactor that is both

environmentally benign and capable of producing high quantities which is a

prerequisite for use in applications.

A dual-stage hydrothermal flow reactor was developed for this purpose and

used to produce hybrid nanomaterials of differing morphologies: from smaller

particles grafted on a larger particle for support, to spherical core-shell

nanoparticles of 10-30 nm in diameter of a narrow size distribution. This is

accomplished by synthesizing the core and shell in separate reactor zones to

avoid a mixed product of ungrafted particles. Using TiO2@SnO2 as an

example the parameters responsible for composite formation are explored.

The used characterization techniques include X-ray diffraction, electron

microscopy and spectroscopy. Structural analysis is applied to demonstrate

the changes in crystal growth mechanism when a material is precipitated

either as a nanoshell or as a free, ungrafted particle. The presence of core-

particles has profound implications for the crystal size, structure and

composition of the nucleating shell material as demonstrated by depositing 2

nm shells of CuO/Cu2O, NiOxLy and Zn6(OH)6(CO3)2 on a 20 nm magnetic core.

SiO2 nanoshells on γ-Fe2O3, TiO2 and α-Fe2O3 cores demonstrate that

morphology depends strongly on the synthesis pH. This factor governs if the

shell is thin and uniform or thick and irregular. An electrostatic model is

developed for identifying synthesis parameters which allows control of the

nanohybrid morphology. The choice of reactants and temperature profile is

shown also to influence whether a composite or mixture is formed.

Multiferroic γ-Fe2O3@SrTiO3 is synthesized to demonstrate the feasibility of

producing highly advanced nanomaterials in flow. The hybrid is obtained

vi

through a complex synthesis route in which the formation of a γ-Fe2O3@TiO2

intermediate is central.

Finally, it is shown how metal sulphides can be produced hydrothermally

using MoS2 as an example. Hydrothermal conditions generally favor oxides,

but choice of precursor can be used to promote formation of polycrystalline

MoS2 nanosheets. The feasibility of producing supported nanomaterials is

demonstrated using TiO2 nanocatalysts on MoS2 sheets

This dissertation describes the dual-stage hydrothermal flow reactor and how

advanced nanocomposites in high yields may be readily synthesized for

potential use in applications.

vii

Dansk resumé

Nanokompositter er en materialegruppe af voksende videnskabelig interesse.

Sammenlægning af to forskellige materialer i en enkelt hybrid nanopartikel

kan skabe multifunktionelle materialer eller bruges til at forstærke

eksisterende egenskaber gennem synergetiske vekselvirkninger. Disse

moderne materialer kan syntetiseres hydrotermalt i en dual-stage flow

reaktor, som er både miljøvenlig og i stand til at fremstille store mængder,

hvilket er en forudsætning for praktiske anvendelser.

En dual-stage hydrotermal flow reaktor blev udviklet til dette formål og brugt

til at fremstille hybrid nanomaterialer af varierende morfologier: fra små

partikler deponeret på en større partikel som support til sfæriske core-shell

nanopartikler 10-30 nm i diameter med en smal størrelsesfordeling. Dette

opnås ved at syntetisere det indre og det ydre materiale i hver sin reaktor for

at undgå et blandet produkt af usammenhængende partikler. Med

TiO2@SnO2 som eksempel udforskes synteseparametrene, der er årsag til

kompositdannelsen

De brugte karakteriseringsmetoder inkluderer røntgendiffraktion,

elektronmikroskopi og spektroskopi. Strukturel analyse er anvendt for at

demonstrere ændringer i mekanismen for krystalvækst, når et materiale

udfælder enten som nanoskal på en partikel eller som en separat

nanopartikel. Tilstedeværelsen af partikler har vidtrækkende konsekvenser

for krystalstørrelse og struktur samt den kemiske komposition for det

udfældede materiale, som demonstreret ved deponering af 2 nm skaller af

CuO/Cu2O, NiOxLy og Zn6(OH)6(CO3)2 på magnetiske nanopartikeler.

SiO2 nanoskaller på partikler af γ-Fe2O3, TiO2 and α-Fe2O3 demonstrer at

morfologien er stærkt afhængig af syntesens pH. Denne faktor bestemmer om

skallen er smal og ensartet eller bred og uregelmæssig. En elektrostatisk

model blev udviklet til bestemmelse af syntese parametre, som giver bedre

kontrol over nanohybridernes morfologi. Valget af reagenser og temperatur

viii

profil er påvist som yderligere faktorer, der bestemmer om nanokompositter

eller en blanding dannes.

Multiferroiske γ-Fe2O3@SrTiO3 er syntetiseret for at vise muligheden for at

fremstille yderset avancerede nanomaterialer i flow. Hybriden blev dannet

gennem en kompleks syntese vej hvorunder dannelsen af γ-Fe2O3@TiO2 er et

vigtigt mellemprodukt.

Til sidst er det vist, hvordan metalsulfider kan dannes hydrotermalt med

MoS2 som eksempel. Hydrotermale vækstbetingelser danner som regel oxider,

men polykrystallinske MoS2 nanoplader kan dannes ud fra de rette reagenser.

Mulighederne for at deponere nanopartikler på større supportpartikler er

påvist ved fremstilling af TiO2 nanokatalysatorer på MoS2 nanoplader.

Denne afhandling beskriver den hydrotermale dual-stage flow reaktor og

hvordan avancerede nanokompositter kan fremstilles i højt udbytte til

praktiske formål

ix

List of abbreviations

Ac Acetate

AIC Ammonium Iron Citrate

ATM Ammonium Tetrathiomolybdate(IV)

BTO Bariumtitanate, BaTiO3

CFD Computational Fluid Dynamics

EDX Energy Dispersive X-ray (spectroscopy)

FTIR Fourier Transform Infrared (spectroscopy)

GUI Graphical User Interface

HPLC High Performance Liquid Chromatography

HAADF High Angle Annular Dark Field

ICP-OES Inductively Coupled Plasma - Optical Emission Spectroscopy

ICSD Inorganic Crystal Structure Database

IEP Isoelectric Point

IPA Isopropyl Alcohol (Isopropanol)

MP Medium Pressure (up to 1500 bar)

p Pressure

PID Proportional, Integral, Derivative

PLC Programmable Logic Controller

PRV Proportional Relief Valve

PXRD Powder X-Ray Diffraction

SS316 Stainless Steel 316

STEM Scanning Transmission Electron Microscopy

T Temperature

TALD Titanium bis(Ammonium Lactato) Dihydroxide

TEM Transmission Electron Microscopy

TEOS Tetraethyl Orthosilicate

TTIP Titanium tetraisopropoxide

XRF X-Ray Fluorescence (spectroscopy)

x

Contents

Chapter 1 Introduction .............................................................................................................. 1 1.1 Materials ............................................................................................................................... 2 1.2 Properties of Nanocomposites ............................................................................................. 3

1.3 Classes of Inorganic Nanocomposites ................................................................................. 4

1.4 Final remarks ....................................................................................................................... 5 Chapter 2 Hydrothermal Conditions ..................................................................................... 7

2.1 The Supercritical Phase ....................................................................................................... 8

2.2 Synthesis of Nanomaterials ............................................................................................... 11 2.3 Industrial Applications ....................................................................................................... 15

2.4 Sub-conclusion .................................................................................................................... 15

Chapter 3 The Dual-Stage Flow Reactor ............................................................................. 17 3.1 Batch and flow limitations ................................................................................................. 18

3.2 Dual-Stage Design .............................................................................................................. 19 3.3 Construction ....................................................................................................................... 23

3.3.1 Hot sections ................................................................................................................. 26 3.3.2 Coolers ......................................................................................................................... 27

3.3.3 Entries and Exits ........................................................................................................ 28 3.3.4 Pump Controller .......................................................................................................... 30 3.3.5 System Monitoring ...................................................................................................... 30

3.3.6 Flow Rates ................................................................................................................... 32 3.3.7 LabVIEW ..................................................................................................................... 33

3.3.8 Safety ........................................................................................................................... 34 3.4 Sub-conclusion .................................................................................................................... 35

Chapter 4 Characterization Techniques ............................................................................. 37 4.1 X-Ray Diffraction................................................................................................................ 38

4.1.1 Generation of X-Rays .................................................................................................. 38 4.1.2 Bragg’s Law of Diffraction .......................................................................................... 39

4.1.3 Structural Analysis ..................................................................................................... 41 4.2 X-Ray Fluorescence Spectroscopy ..................................................................................... 42

4.3 Scanning Transmission Electron Microscopy ................................................................... 43 4.4 Fourier Transform Infrared Spectroscopy ........................................................................ 44 4.5 Sub-conclusion .................................................................................................................... 44

Chapter 5 Commissioning ....................................................................................................... 45 5.1 TiO2 ..................................................................................................................................... 46

5.2 Flow synthesis of TiO2 ....................................................................................................... 47

5.3 Sub-conclusion .................................................................................................................... 51 Chapter 6 Formation of a TiO2@SnO2 Nanocomposite .................................................... 53

xi

6.1 Temperature and pH .......................................................................................................... 54 6.2 Composite formation .......................................................................................................... 62 6.3 Concentration Effects ......................................................................................................... 64

6.4 Evaluation of Secondary Hot Mixing................................................................................. 66 6.5 Mechano-chemical processing ............................................................................................ 67 6.6 Sn-Precursor ....................................................................................................................... 68 6.7 Sub-conclusion .................................................................................................................... 71

Chapter 7 Silica .......................................................................................................................... 73 7.1 SiO2 Shells ........................................................................................................................... 74

7.1.1 γ-Fe2O3@SiO2 ............................................................................................................... 76 7.1.2 TiO2@SiO2 .................................................................................................................... 78 7.1.3 α-Fe2O3@SiO2 ............................................................................................................... 80

7.2 Sub-conclusion .................................................................................................................... 83

Chapter 8 IEP and Morphology ............................................................................................. 85 8.1 IEP predictions ................................................................................................................... 86 8.2 Copper-based Shell ............................................................................................................. 88

8.3 Nickel-based Shell .............................................................................................................. 91 8.4 ZnO and Zn6(OH)6(CO3)2 .................................................................................................... 95

8.5 Sub-conclusion .................................................................................................................... 99 Chapter 9 γ-Fe2O3@TiO2 ......................................................................................................... 101

9.1 Magnetic photocatalysts ................................................................................................... 102

9.2 The influence of precursor and concentration ................................................................ 102 9.3 Sub-conclusion .................................................................................................................. 107

Chapter 10 Multiferroics ....................................................................................................... 109 10.1 BaTiO3 and SrTiO3 Ferroelectrics ................................................................................. 110

10.2 Reactor Reconfiguration ................................................................................................. 112 10.3 BaTiO3 Precursors .......................................................................................................... 113

10.3.1 Ba(OH)2 .................................................................................................................... 113

10.3.2 Ba(NO3)2 ................................................................................................................... 116 10.4 The Two Synthesis Steps ............................................................................................... 117

10.4.1 Step 1: Titania deposition ....................................................................................... 118 10.4.2 Step 2: Conversion of Titania to Sr- and Ba-Titanate .......................................... 119

10.5 Precipitation of carbonates ............................................................................................ 123

10.6 Evaluation of synthesis strategy ................................................................................... 125

10.7 Effect of pH...................................................................................................................... 127 10.8 Sub-conclusion ................................................................................................................ 129

Chapter 11 MoS2 ....................................................................................................................... 131 11.1 Structure and applications ............................................................................................. 132

11.2 Hydrothermal synthesis of phase pure MoS2 ............................................................... 132 11.3 TiO2 particles on MoS2 substrates ................................................................................. 137 11.4 Sub-conclusion ................................................................................................................ 142

Chapter 12 Concluding Remarks ......................................................................................... 145 References ................................................................................................................................. 147 Appendix ..................................................................................................................................... A1

1

Chapter 1

Introduction

This dissertation describes the design, construction and performance of a

novel dual-stage flow reactor for hydrothermal synthesis of advanced hybrid

nanomaterials. The performance of nanocomposites are of great interest for

catalysis, energy storage and many other applications which form the

motivation for a green, and sustainable high-yield synthesis route to produce

them. Chapter 1 describes the new and interesting properties that can arise

when different materials are combined into a composite on the nanoscale.

Many different combinations exist, and this chapter describes some of the

most important classes of inorganic nanocomposites.

Chapter 1 Introduction

2

1.1 Materials

The history of humankind is a history of materials. A story of how we have

used the resources found in our surroundings and shaped them into materials

that we could use in our daily lives. Our development has been defined by the

resources we learned to extract, and the materials we were able to produce. In

the Stone-age rocks were collected directly off the ground and shaped into

tools but it was in the Bronze-age that we learned that rocks could be heated

in the fire and molten metals could be extracted. Using a furnace instead of a

simple bonfire resulted in higher temperatures and allowed for extraction of

iron which started the Iron-age, and later the addition of carbon created steel.

The use of ceramics allowed us to construct skyscrapers of concrete, plastics

showed us how artificial materials not found in nature can be synthesized,

and as such materials have guided our development into our current age

which is often referred to as the Silicon-age.

It is the properties that make a material. Material properties are intimately

linked to the structure and are the reason why diamond is hard and graphite

is not. Nanoparticles are an interesting group of materials in this respect as

their properties often differ greatly from the bulk. Moreover, the low particle

size results in a material with a great surface area for chemical reactions to

occur. The structure of nanomaterials can be determined through the

synthesis procedure by how the building blocks are arranged during

formation. It is therefore the initial synthesis that ultimately decides the

properties of the material and its performance in applications.

When two different materials come into contact they may interact with each

other. For nanoparticles these interactions can be very strong because of the

great contact area. Such nanocomposites are highly interesting materials

because of the novel properties that may emerge, but they are inherently

difficult to synthesize. Synthesis methods that offer great control of the

deposition of one material on top of another are usually incapable of

producing the high yields necessary for use in applications such as batteries

and fuel cells. This dissertation demonstrates how nanocomposites can be

1.2 Properties of Nanocomposites

3

synthesized hydrothermally and describes a reactor type that can be used to

perform this in a continuous flow.

1.2 Properties of Nanocomposites

The combination of two or more materials into a single material offers a way

to not only obtain a multifunctional material but also enhance existing

properties or even result in new properties. Such effects arise from the

interactions between the materials along their interface at which they form a

physical contact. Nanoparticles have intrinsically large surface/volume ratios,

and nanoparticles with a coating of a different material are ideal for

optimizing the contact area between two materials. Such core-shell

nanoparticles may produce synergistic effects i.e. properties greater than the

sum of the individual components, and this can be used in applications.

In catalysis the performance of fuel cells has been enhanced by deposition of

separate layers of gold and platinum on iron oxide. The electrocatalytic

activity of the hybrid was found to surpass that of the individual constituents

for the methanol oxidation reaction and oxygen reduction reaction.1 The

results also provide an example of how gold and other rare and expensive

materials can be deposited as coatings on a cheaper core in order to obtain a

composite particle which largely has the appearance of a large gold particle.

The optical properties of CdSe can be tuned by deposition of a CdS shell

around the semiconductor core. The large lattice mismatch between core and

shell induces a strain that persists for several nanometers across the

interface which can be used to change the absorption spectra for use in light-

emitting-devices.2, 3

Magnetic properties can be enhanced by exchange coupling between the hard-

magnetic core and soft-magnet shell through direct contact. The maximum

energy-product of the FePt@Fe3O4 nanocomposite was shown to be 38 %

higher than that of single-phase FePt.4

Chapter 1 Introduction

4

Applying a core-shell structure was shown to improve both the cyclability and

the electrochemical performance of Li-ion batteries. The role of the shell

material is numerous including providing a buffer space for the expansion

and contraction of the electrode material due to lithium ion migration during

charge and recharge cycles, improving the electronic conductivity and

protecting the active material which may be air sensitive.5

1.3 Classes of Inorganic Nanocomposites

Many nanocomposites consist of multiple inorganic nanoparticles grafted on a

single larger particle for support. The role of the support is primarily to fixate

the nanomaterial but can also enhance the performance of the catalytic

material. An example is carbon flakes which are good electronic conductors

and have been used as support for platinum nanoparticles in fuel cells6.

Inorganic materials such as ZnO7, TiO28, Al2O3

9, 10 and SiO211 have also been

employed as support materials. If the secondary coating forms a closed layer

the support material is then referred to as the core while the coating is the

shell. The term core-shell is, however, often used generically in literature and

does not always indicate a closed shell of uniform thickness surrounding a

near-spherical core. In this dissertation the term core-shell is used mainly as

a morphological description of the nanocomposite but also to better facilitate

the results to the reader.

Another class of hybrid materials is inorganic nanoparticles functionalized by

deposition of organic molecules. The role of the organic shell can be to

improve dispersibility or to protect the core from oxidation, but it is especially

interesting for biological applications as it allows the inorganic nanoparticle

to interact with the cells.12 Organic shells can also be produced in-situ from

the precursor or a surfactant.13, 14

For these reasons hybrid materials are highly interesting for use in

applications.15 However, syntheses of nanocomposites are generally

associated with great difficulty and the multistep synthesis procedure is

1.4 Final remarks

5

labour intensive, and often only minute quantities are produced. While this

may suffice for characterization purposes gram-scale synthesis is necessary

for applications testing.

1.4 Final remarks

To address these issues a dual-stage hydrothermal flow reactor was

constructed which combines great control of particle characteristics with a

high production capacity. The dual-stage is the latest generation of

hydrothermal continuous flow reactors and consists of two connected flow

reactors for separate formation of the primary and the secondary phase. With

this powerful synthesis tool a range of nanocomposites has been synthesized

to demonstrate a green and sustainable method for production of advanced

nanomaterials, and this dissertation aims to understand the synthesis

parameters responsible for composite formation.

7

Chapter 2

Hydrothermal Conditions

At elevated pressure and temperature the characteristics of water change

which has a remarkable effect on its properties. This chapter discusses these

changes and how they can be used to tailor nanomaterials. The principles

behind hydrothermal flow synthesis are described, and an overview is

provided of some of the research groups and companies which use

hydrothermal conditions for a green and sustainable synthesis of

nanoparticles. The chapter ends with how supercritical solvents are used in

applications outside of materials science.

Chapter 2 Hydrothermal Conditions

8

2.1 The Supercritical Phase

As water is heated the vapour pressure of the liquid will rise until it is equal

to the external pressure, and the water boils. If the liquid is pressurized the

boiling point will shift to higher temperatures. The temperature (T) and

pressure (p) dependency of the boiling point is known as the liquid-vapour

phase boundary, and a plot of density of water and steam as a function of T

and p is shown in Figure 2.1. The phase transition is located along the density

discontinuity, but at the critical point (374°C and 221 bar) the discontinuity

disappears. The surface boundary once separating the liquid and the gas has

collapsed, and the fluid has entered a single-phase state. This phase is known

as the supercritical phase, and the region is highlighted in Figure 2.1. Here

the net attractive and repulsive intermolecular forces are comparable in

magnitude, and a thermodynamic driving force towards phase separation no

longer exists.

Figure 2.1 | Density of water as function of pressure and temperature with supercritical

region highlighted. Note the inverted temperature axis. Plotted using Density function in

IAPWS_IF97 MATLAB package16 based on the International Association of Water and

Steam Industrial Formulation 1997.17

2.1 The Supercritical Phase

9

The solvent properties of supercritical fluids can be tuned because the density

changes continuously. Gas and liquid densities are separated by 3 orders of

magnitude and as the kinetic and potential energies of the water molecules

become equal the solvent properties become intermediate to those of the gas

and the liquid. The liquid-like densities result in high momentum transfers

while transport properties such diffusivity and viscosity are more gas-like.

Water is a highly polar solvent due to its ability to form hydrogen bonds, and

these weak interactions are indeed what set water apart from the other

hydrides of group VI. A decrease in density causes the average distance

between water molecules to increase. This reduces the number of hydrogen

bonds each molecule is able to form to its neighbors, and this has a profound

effect on the dielectric constant (Figure 2.2) which is a measure of the polarity

of the solvent. It is 80 at room temperature but drops to 2 at 400°C which is

comparable to non-polar organic solvents such as hexane. Ionic species in

solution are unable to remain solvated in low polarity solvents.18

Many chemical processes are pH dependent and hydrothermal conditions

may enhance acid/base catalyzed reactions without the need for added acids

or bases. The dissociation constant, Kw, of water (Figure 2.3) changes by

several orders of magnitude as the temperature increases, and the

equilibrium

2𝐻2𝑂 ↔ 𝐻3𝑂+ + 𝑂𝐻−

shifts to the right. The ionization product of water is given by

𝐾𝑤 = [𝐻3𝑂+][𝑂𝐻−]

Kw is 10-14 at ambient condition but reaches a maximum of 10-11 around 250°C.

With higher temperatures Kw drops due to the decrease in dielectric constant

which destabilizes solvation of ionic species.


10

Figure 2.2 | Dielectric constant, ε, of water and steam. Water becomes less polar at higher

temperatures. Plotted using IAPWS_IF9716 and Dielectric function by Uematsu and

Franck.19 Supercritical region is highlighted.

Figure 2.3 | The ionization product of water reaches a maximum around 250°C. Plotted

using IAPWS_IF9716 and Ionization product function supplied by Marshall and Franck.20

2.2 Synthesis of Nanomaterials

11


The term “hydrothermal” was first used in the field of geology and described

the precipitation of rocks and minerals from water at the elevated

temperatures and pressures found in the crust of the Earth. It comes from the

greek words “hydros” and “thermos” which translates into water and heat,

and is used to describe temperatures above room temperature and pressures

above atmospheric pressure.21 For many materials scientists the term is more

specifically used to describe water in the regions closer to the critical point

otherwise referred to as sub- near- and supercritical water.

In 1992 Adschiri et al.22, 23 demonstrated that supercritical conditions can to

be used to obtain metal oxide nanoparticles in a continuous flow setup. Fast

heating of a metal salt resulted in a burst of nucleation, and the

nanoparticles produced were of a very narrow size distribution, because of the

quick supersaturation. The low solubility and high reaction kinetics meant a

rapid depletion of the precursor with little to no material available for further

growth of the nanoparticles. In addition, the short reaction time meant a

suppression of further particle growth.

The formation of a crystalline solid is often a complex process and detailed

knowledge of reaction mechanisms have been provided by in-situ studies.24, 25

However, a generalized reaction pathway from a metal salt to metal oxide can

often be considered as consisting of two main reactions: hydrolysis of the salt

to form the hydrous oxide, followed by dehydration.

Hydrolysis MLx + xOH−

→ M(OH)x + xL−

Dehydration M(OH)x

→ MOx2

+ x

2H2O

In a typical batch synthesis at hydrothermal conditions the slow heating rate

will allow hydrous oxides to grow into large particles before dehydration sets

in. By improving the heating rate dehydration can be achieved before


12

hydrolysis is completed, and the resultant particle size is drastically reduced.

This is further aided by an increase in dehydration rate for smaller particles.

A rapid heating of the precursor is key, and this was done by mixing the

precursor with superheated water in a continuous flow apparatus and leading

the contents into a tubular reactor. The hot solvent, precursor and reactor

tubes were connected in a tee-fitting, and this mixing geometry is therefore

known as tee-mixing. Since 1992 supercritical flow reactors have continued to

follow the general outline of the original reactor design by Adschiri and co-

workers. A significant change in mixing geometry was the introduction of the

Nozzle mixer developed by Lester and co-workers in Nottingham.26 This setup

uses a counter-current mixing geometry and has been used for the synthesis

of several nanomaterials.27-31 Mixing is improved by the radial symmetry as

calculated by Computational Fluid Dynamics (CFD) and illustrated using in-

situ diffraction tomography.32-34

In Bordeaux supercritical synthesis has been coupled with microfluidics by

Aymonier and co-workers. Water and alcohols are frequently used solvents,

but carbon dioxide and ammonia has been demonstrated as useful solvents

with the latter useful for metal nitrides. Hybrid nanomaterials have been

synthesized by the deposition of organic molecules onto inorganic

nanoparticles.35-41

Micro tee-mixers of reduced inner diameter were investigated by Kawasaki

and co-workers. Reductions in inner diameter and increased flow rates were

shown to decrease the nanoparticle size as the improved mixing lead to an

enhancement in the nucleation rate. Swirl-mixers are another type of mixer

in which the precursor is injected into the middle of a vortex created by two

“swirling” supercritical streams which has also been demonstrated as a

powerful synthesis tool, although the mixer must be custom-made.42-46

A tee-mixer on a vertical reactor allows for three different configurations by

changing the relative orientation of the precursor and the solvent. These all

have different mixing behavior which changes with flow rates and with the


13

dimensions of the tubing and mixing tee. A high solvent flow rate will provide

a faster heating, but due to density differences the mixing volume is

increased as the hot solvent flows penetrates deeper into the precursor inlet.

The resulting size distribution is therefore widened with higher flow rates.

Conversely, a rapid heating is necessary for a fast nucleation step. CFD

calculations and neutron imaging have recently been compared with product

size distributions for hydrothermal synthesis of CeO2 nanoparticles.47 In

general, a solvent flow rate which is two to three times larger than the

precursor gives the best results in terms of low average size and narrow size

distribution.

Supercritical synthesis is commercially interesting as supercritical water is

both cheap and environmentally benign. Changes in solvation typically

require organic solvents, but the tunable properties of supercritical water

means that synthesis can be performed without the use of expensive and/or

toxic chemicals. The waste product from supercritical synthesis is primarily

water, and supercritical water is therefore considered a green chemical. South

Korean Hanwha Chemical Corporation was the first company to

commercialize the technology in 2011. Current production capacity is around

150 kg/h of LiFePO4.48 In Nottingham Promethean Particles is driving the

technology towards commercialization while in Denmark the push is done by

Teknologisk Institut. In Japan a large-scale industrial synthesis facility is

currently under construction.49

In 2008 a hydrothermal flow reactor was constructed at Department of

Chemistry in Aarhus by Peter Hald based on the original design by Adschiri.

A diagram of the reactor is illustrated in Figure 2.4.

Superheated solvent and precursor are pumped into the system by pumps

placed on either side of the reactor. The two streams are mixed and led into a

vertical heating zone for further particle growth and crystallization. The

vertical heating zone is the reactor but this term is often used

interchangeably with the entire flow system consisting of pumps, heaters,

reactor zone, valves etc.


14

Figure 2.4 | Outline of single-stage hydrothermal flow reactor constructed by Peter Hald.

Precursors are loaded into an injector, and the content is displaced towards the reactor

zone for nanoparticle formation. Mixing temperature is measured by a thermocouple (TC).

After maturing inside the reactor the suspension is cooled down, and the

product is collected downstream of a proportional relief valve (PRV) which

both acts as reactor exit and pressure control. For many syntheses the

superheated solvent not only provides a fast heating rate but is in fact also a

reactant e.g. as oxygen-donor. Although a simplification, the heated stream

will be referred throughout the text as solvent stream.

A significant addition to the Aarhus flow reactor was the installment of a

custom-made injector. A common problem in flow synthesis is pump failure

due to corrosion, clogging from viscous gels or sedimentation from unstable

suspensions. These problems arise because the moving parts of the pump are

in direct contact with the precursor, and this was solved by the use of an

injector. It consists of a 200 mL tube with a movable piston that separates

precursor and the pumping solvent - usually water or an alcohol. This made

the reactor more versatile but also laborious and the semi-continuous flow is

unsuitable for extended production runs. Additionally, an interrupted flow is

not a viable solution for any downstream treatments such as mixing with a

2.3 Industrial Applications

15

secondary precursor for nanocomposite formation which requires a steady

supply of particles for extended periods. These problems had to be addressed

and were solved by the choice of design and equipment for the dual-stage

reactor as discussed in the following chapter.

2.3 Industrial Applications

Supercritical fluids are highly useful for inorganic synthesis, but they are also

used in several areas outside of this field. The high miscibility of fuel and

oxidants combined with the high reaction rates present in a single phase

medium have been utilized for the destruction of organic wastes for decades.

This process is known as supercritical water oxidation. Plants are under

construction around the world for the destruction of sewage and chemical

weapons such as the Blue Grass Chemical Agent-Destruction Pilot Plant in

Richmond, Kentucky by General Atomics.50

In the food industry SCFs have been used to extract vitamins, oils, caffeine,

cholesterol and other compounds.51 The largest industrial application,

however, is as a coolant in fossil fuel and nuclear power stations due to the

high heat capacity near the critical point.52

2.4 Sub-conclusion

The properties of water change drastically with increases in temperature and

pressure. The drop in density forces the water molecules further and further

apart which has profound implications for its behavior. The changes have

been used to synthesize advanced nanomaterials. Hydrothermal flow reactors

are of particular interest as they provide a fast heating, and the nanoparticles

produced are very small and of narrow size distribution.

Hydrothermal flow reactors are of rising commercial interest because it is a

scalable technology and the main chemical, water, is cheap. What may be


16

otherwise accomplished using expensive and/or toxic chemicals can be

performed using only water and its behavior at hydrothermal conditions. The

synthesis method is therefore very environmentally benign, and supercritical

water is considered a green and sustainable chemical.

Hydrothermal flow reactors are useful for producing single-phase

nanomaterials but not for nanocomposites. A different type of flow reactor is

necessary for such compounds and Chapter 3 discusses the design and

construction of such a reactor type.

17

Chapter 3

The Dual-Stage Flow Reactor

Over the course of the last decades hydrothermal synthesis have been

demonstrated as not only a green and sustainable but also highly versatile

chemical route for producing nanomaterials. This chapter opens with a

description of the synthesis conditions necessary for composite formation and

how choice of design can be used to promote heterogeneous over homogeneous

nucleation.

Flow reactors offer a high degree of product control, but the accelerated

reaction kinetics generally make them unsuitable for hybrid nanomaterials.

Conversely, the slow heating in batch reactors is more suitable for obtaining

such hybrids but at the expense of both product characteristics and quantity.

This forms the motivation for construction of a new type of flow reactor.

The chapter discusses the construction of this new reactor system termed

“The Dual-Stage Flow Reactor”. The chapter aims at covering all aspects of

the reactor engineering; from choice of design and equipment to installation

and software programming. The performance of the reactor system was

evaluated through commissioning with TiO2 which is the subject of Chapter 5.

Chapter 3 The Dual-Stage Flow Reactor

18

3.1 Batch and flow limitations

Nanoparticles in suspension are denser than their surrounding medium, but

because of their large surface area they can remain suspended through

interactions with the medium. The nanoparticles and precursor can therefore

form a stable mixture, although the particles may sediment over time. Batch

synthesis of nanoparticles is by far the most common synthesis method due to

its simplicity in terms of precursor and product handling and the equipment

required. Synthesis of composite nanoparticles in batch is usually performed

by a two-step process. First, the core nanoparticles are synthesized and

secondly the shell precursor is added to the reaction vessel.

For synthesis at ambient condition the secondary precursor is simply added

to the glass beaker once the core nanoparticles are fully formed and dispersed.

For hydrothermal synthesis, however, the procedure involves the use of a

steel autoclave which has to be cooled down prior to addition of the secondary

precursor. The low reactor volume and laborious process significantly reduce

the obtainable yield. Furthermore, the slow heating (more than one hour to

reach 90% of the set-temperature) will lead to sedimentation and aggregation

of the core particles. As the secondary precursor nucleates this will lead to

partial homogeneous nucleation of the secondary material, and the product

will be a physical mixture of uncoated core material, ungrafted shell material

and single particle nanocomposites of differing core/shell volume ratios.

Both core and shell precursors may in principle be loaded into the same

reactor vessel but only if the shell nucleates at a higher temperature than the

core nanoparticles, and provided that no intermediate phases or solid

solutions are formed and that the core particles are fully suspended at the

onset of secondary nucleation. This is rarely the case.

Heating two precursors simultaneously often lead to intermediate phases.53, 54

Conventional single-stage flow reactors are therefore unsuitable for

nanocomposite syntheses as nucleation and growth of both phases are

confined to the same mixing piece and reactor volume. Synthesis of

3.2 Dual-Stage Design

19

nanocomposites in single-stage reactors could be performed by mixing the

precursor with a premade dispersion of core nanoparticles and pumping the

mixture into the reactor. This, however, doubles the synthesis time as the

core material must first be synthesized, and the strain on the pumps is

significant.

The latter problem can, however, be solved by the use of injectors; the

suspension and precursor is mixed and loaded into a tubular container,

placed downstream of a pump (Figure 2.4), and the injector content is

displaced towards the reactor zone by the incoming solvent. Double-injector

pumps for continuous flow are commercially available (e.g. “Syringe Pump”,

Teledyne Isco). Nevertheless, several problems may arise including

sedimentation of particles, increased user workload (core-particles must be

pre-synthesized and injectors must be continuously refilled) and general

operating problems (erosion, clogging etc.) associated with particle-carrying

streams in contact with moving mechanical parts. A diluted dispersion will be

stable towards sedimentation, but this reduces the yield by orders of

magnitude compared to using a concentrated (and unstable) suspension.


The outline of the dual-stage flow reactor is reported in Publication I. It was

designed to mimic the two-step synthesis procedure from batch methods

while overcoming the limitations of single-stage flow reactors by dividing

formation of core and shell materials into separate reactor zones. The general

synthesis outline is the formation of a primary nanoparticle suspension which

is cooled down to near room temperature, mixed with a secondary precursor,

and the resulting mixture then led into a secondary reactor zone for

formation of the secondary material. The use of multiple reaction zones

connected in series has previously been used in microfluidics.55 A diagram of

the dual-stage reactor system is shown in Figure 3.1, but the modular design

allows for several different configurations.


20

Figure 3.1 |The shown dual-stage configuration was employed for synthesis of a

TiO2@SnO2 nanocomposite from TTIP and SnCl4. Heated zones are red, and cooled zones

are blue.

A more direct approach would have been to simply cut the vertical reactor

zone of a conventional single-stage flow reactor in half, insert a tee-mixer and

inject fresh precursor straight into the hot reactor contents. This is a much

simpler method but was deliberately discarded. By choice of design the dual-

stage hydrothermal reactor offers five distinct advantages over what will be

referred to as the “direct approach”.


21

1. A smaller temperature gradient

Homogeneous nucleation requires several atoms to come into proximity and

overcome the high surface energy of the unstable nuclei. The energy barrier

for precipitation can be lowered if a stable surface is already present and the

precursor will nucleate heterogeneously on this pre-existing surface.

Deposition of a shell material is therefore expected to take place at a lower

temperature than that required for homogeneous nucleation. The flow rate of

the cold stream feeding into the secondary reactor is the sum of all flow rates

used in the primary reactor and the secondary precursor streams. This

increased mass flow is heated more slowly than in the primary reactor and

will promote heterogeneous nucleation as the precursor is exhausted before

the temperature is high enough for homogeneous nucleation.

2. No premature heating of secondary precursor inlet

In the direct approach the lower density of the hot reactor contents would

spill into the horizontal precursor inlet. This causes the inlet to be

prematurely heated which will promote homogenous nucleation,

sedimentation and eventual blockage. Contrarily, in the dual-stage design the

product of the primary reactor is cooled down prior to mixing with the

secondary precursor. The secondary precursor will therefore not precipitate in

most cases, but this also depends on the chemicals used.

3. Mixing prior to reactor entry

In the direct approach the shell material is simultaneously mixed with and

heated by the nanoparticle suspension. Asymmetric concentration and

temperature gradients are then formed across the reactor tube downstream

from the inlet. Mixing will not be completed by the time of secondary material

formation, and the product is a mixture of uncoated core nanoparticles,

ungrafted shell material, and nanocomposites of varying core/shell volume.


22

This is not a problem in the dual-stage reactor because mixing is performed

outside the heated zones.

4. Independent secondary reaction temperature

The direct approach offers no temperature control of the shell synthesis. This

is particularly problematic for materials such as SiO2 or ZnO which

precipitate at “low” reaction temperatures (50-100°C). In a typical

hydrothermal synthesis a precursor is pumped at 5 mL/min and mixed with

superheated solvent at 15 mL/min and then heated in the reactor at 350°C. If

this hot suspension (20 mL/min) of core nanoparticles is mixed with a SiO2

precursors pumped at 5 mL/min at room temperature it would lead to

immediate and uncontrolled shell formation. A gradual increase in

temperature is necessary in order to reach the heterogeneous nucleation

temperature prior to the homogenous, and this cannot be done by the direct

approach.

5. Reactor length

The combined flow rate in the secondary reactor is higher than in the primary

reactor, which decreases the residence time (both reactors are of identical

length and diameter). Conversely, a slow heating of the reaction mixture

necessitates an increased residence time. The length of the secondary reactor

tube should therefore be maximized.

In the direct approach the second reactor is placed directly below the first

reactor. The length of each of these two reactors is therefore half that of

single-stage reactor. Placing the primary and secondary reactors vertically on

top of each other would require a laboratory of unusual ceiling height without

compromising cooler performance and reactor length.

3.3 Construction

23

By cooling down the primary nanoparticles the dual-stage reactor design

enables the suspension to be transported horizontally and upwards into the

top of the secondary reactor. This allows the reactor length of the secondary

reactor to be extended to above one meter and typical residence times are of

the order of a minute for each reactor stage. Horizontal reactor zones are

generally avoided as particles tend to sediment inside the reactor tube.

3.3 Construction

Mounting both reactor stages next to each other in the same plane would

require a very broad stretch of wall not to mention several tubes having to

cross over. The solution was to “fold” the entire reactor setup around a grid

wall with the primary reactor stage on the “front” side and the secondary

reactor on the “rear” side as its mirror image. This drastically reduced the

dead-volume from the long sections of tubing otherwise required to transport

the contents. Installation of controllers, pumps and precursor containers at

opposing ends of the grid wall allows the user unhindered access to both

reactor stages.

Figure 3.2 and Figure 3.3 show the front and rear side of the reactor setup.

Prior to construction the setup was designed and visualized by modelling in

Autodesk Inventor and a 3D rendering of the final model can be found in

Appendix 1. Note that solvent heaters were installed for both reactor stages

to investigate the influence of heating rates in the secondary reactor.

A high flow speed is necessary to prevent sedimentation in the horizontal

streams and gravitational drag in the upwards going streams carrying

particles. However, this cannot be done by simply increasing the flow rate as

this would have profound effects on the formation chemistry, residence time,

heat exchange etc. The solution was therefore to reduce the tube diameter

which had the added benefit of further decreasing the system dead-volume.


24

Figure 3.2 |Front side of dual-stage showing the primary reactor. Flow rates are regulated

by black controller box (bottom left), and pressure is monitored from gauges. Cooling water

flow rates are controlled for each main cooler.

3.3 Construction

25

Figure 3.3 |Rear-side of dual-stage. The secondary reactor is suspended on a grid wall to

the right. Precursors are pumped from plastic bottles by the three yellow Milroyal D

pumps. Temperatures are regulated by PID controllers placed in control boxes (top left).


26

Medium pressure (MP) tubing of 1/4 inch’’ (2.8 mm I.D.) tubing was used for

most of the cold tubing sections, because this is the smallest diameter

available for MP tubing. MP tubing was selected as it enables virtually

unlimited assembly and disassembly which is necessary for cleaning and

configuration changes.

3.3.1 Hot sections

Heaters for the superheated solvent stream and for the reactor zones were

constructed by insertion of heat cartridges into aluminium for low weight and

high thermal conductivity. Each heater (50x50x450 mm) consisted of two

blocks held in close contact with a 3/8’’ tube by four screws as illustrated in

Figure 3.4. Ten 100 W heat cartridges were used per heater. A threaded

thermocouple was screwed into each heater for better contact and improved

heat transfer. Each thermocouple was connected to a Eurotherm 3216 PID-

controller for temperature control. This unit opens and closes the heater

circuit in response to the difference in measured and setpoint temperature.

The proportional, integral and derivative terms in the algorithm ensures that

the temperature does not over-shoot or fluctuate but stabilizes quickly.

Figure 3.4 | Sketch of a main heater. The heat cartridges and thermocouples are

connected to the PID temperature controller.

3.3 Construction

27

For the superheated solvent Inconel625 tubing was used due to its enhanced

corrosion resistance at supercritical conditions. The reactor tube was

constructed from SS316. Stainless steels are less resistant to corrosion but

much cheaper. This enables the user to switch reactor tube and reduce cross-

synthesis contamination. At the end of its service life a tube is simply

discarded and replaced. A dozen reactor tubes have been constructed for this

purpose. All MP tubes were purchased from BuTech while fittings and valves

were purchased from Swagelok.

Two heaters were used for the approximately 1100 mm long reactor tube and

another two were used for the solvent stream. In addition, a third solvent

heater was installed on the solvent stream prior to the main aluminium

heaters. This preheater consisted of a hollow aluminium cylinder with two

300 W heat cartridges inserted from either end with 6 m 1/8’’ Hastelloy (C-

276) tubing coiled around. The preheater was installed to aid heating at

higher flow rates. The single-stage uses up to 4.6 kW, while the dual-stage

setup uses 6.6 or 9.2 kW depending on whether the secondary solvent heaters

are in operation.

3.3.2 Coolers

Tube-in-tube heat exchangers were constructed in-house at the chemical

workshop. The purpose of a cooler is to quench the suspension and stop the

chemical reactions. To collect the product the product temperature must be

below the boiling point otherwise it would exit as a jet of hot steam. However,

the rubber seals in the pressure relief valve (PRV) are rated to a maximum

temperature of 70°C, and for practical purposes it is just more convenient to

handle a suspension that is near room temperature. The 350 mm straight

tube used as cooler effectively cools the reactor contents for most syntheses.

Exit temperature is evaluated manually by touching the steel PRV with the

hand. Steel gets “too hot” to touch above 50°C.


28

Inadequate cooling is observed when combining high flow rates with elevated

reaction temperatures, but is only a real problem when large quantities of gas

are generated as the flow velocity is greatly increased. In case of cooler failure

the hot contents would reach the PRV and damage the rubber seal. The hot

pressurized contents would rapidly expand by the pressure drop and exit the

reactor at an explosive rate.

The purpose of a cooler is also to isolate the heated regions. During operation

heat is transported with the flow down towards the main cooler located

between the vertical reactor and the PRV, but the heat will spread at

stagnant flow in the steel tubing by conduction and warm up the precursor

inlet leading to premature precipitation. For this purpose a small cooler is

installed on the precursor inlet to the mixing tee to reduce conductive heating.

An identical cooler should in principle be placed between the solvent pump

and solvent heaters, but the two are far apart and connected by 1/16’’ tubing

which is rapidly cooled by the ambient air.

3.3.3 Entries and Exits

The precursors and solvents are pressurized and injected into the system by

pumps. Three Milroyal D membrane pumps (Milton Roy) were installed for

use as precursor pumps (93 strokes/min, maximum displacement volume 0.71

mL/stroke). At the first half of a pumping cycle a steel piston exerts a force on

a moving steel membrane inside the pump head. The content is displaced in

the direction towards the reactor as determined by the check valves. At the

final half the piston moves opposite and the generated vacuum draws fresh

liquid inside the pump head. The membrane is the only moving part in

contact with the medium which makes the pumps very robust and capable of

handling even dilute dispersions and gels; thick dispersions and gels tended

to clog the pumps. Pumps are run continuously and three-way valves are

installed to switch between precursor and solvent.

3.3 Construction

29

Double-acting HPLC pumps (LabAlliance Prep 36) were selected for solvents,

as the dual pump head configuration gives a less pulsating flow. These pumps

are less stable and require degassing by ultrasonication under vacuum to

remove air bubbles in the water prior to use. Earlier syntheses were aborted

after a couple of hours in operation due to HPLC pump failure. After

degassing of the demineralized water no pump failure has occurred.

At the reactor exit downstream of the cooler is a PRV from which the product

is collected. During pressure build-up a spring is compressed by the

pressurized medium. The spring is attached to a piston which is finally

pushed back far enough to allow the contents to escape via an opening. This

causes the system pressure to drop, the force exerted on the spring

diminishes, the piston retracts and the opening closes until the pressure has

been rebuilt.

A PRV is not designed for maintaining a stable pressure but to prevent the

pressure from exceeding an upper value. PID-controlled backpressure

regulators are generally used for this purpose. Such a regulator continuously

measures the upstream pressure, and this signal feeds into the PID-controller,

which opens and closes the exit valve in response to the pressure signal. This

provides a much more stable system pressure. Nevertheless, the regulator is

a more expensive solution over the PRV especially when considering

replacement parts due to corrosion and erosion from highly concentrated

suspensions. Typical pressure fluctuations from PRVs are approximately ±5

bars in operation. Pressure fluctuations decrease with temperature as the

medium in the hot zones becomes more compressible, but these pressure

fluctuations increases when particles are present in the fluid. A reactor

system without small variations in pressure has a very stable flow speed but

this will promote sedimentation on the reactor walls. The PRV was therefore

chosen for its simplicity and because they are quick to remove, clean and

reinstall between syntheses.


30

3.3.4 Pump Controller

The three precursor pumps are all fitted with an actuator. This little motor

sets the dispensed volume from an analogue input signal of 4-20 mA. Each

electrical signal is generated by a turn dial potentiometer connected to a

display for an accurate current readout (Figure 3.5). This allows fast and easy

remote control of the pumps instead of having to manually adjust the knobs

on each pump which are not easily accessible. 4 mA and below translates into

0% while 20 mA is 100% pump capacity.

Figure 3.5 | Pump controller box. Flow rate is increased by a turn dial with a LED for

better readout.

3.3.5 System Monitoring

The addition of a secondary reactor and extra inlets meant the reactor system

was to become more complex, and it was evident that an automatic

monitoring of the reactor system was necessary. A Compact RIO (CRIO 9076)

Programmable Logic Controller (PLC) from National Instruments was

purchased to sample data from pressure transmitters (Armaturenbau) and

thermocouples (BuTech) and transmit it to a separate computer operated by

the user. Logging of the data would provide an overview of the reactor state

and synthesis conditions, and a Graphical User Interface (GUI) was therefore

developed. The system monitoring was later expanded with an automated fail

safe protocol to shut down pumps and heaters in case of pressure build-up,

usually due to clogging. A flow chart of the PLC environment is shown in

Figure 3.6, and the GUI is shown in Figure 3.7. These implementations made

3.3 Construction

31

the reactor system much less “black-box” and more transparent and user-

friendly.

Figure 3.6 | Flow chart of PLC environment. PLC samples data from sensors and scales

(via a server). Users access data via the GUI. Power circuit for heaters and pumps are

opened if pressure exceeds 325 bar by an actuated switch (red diamond).

Figure 3.7 | GUI front page provides overview of flow rates and densities (left) and

pressures (right). Charts containing flow, pressure, temperature, settings and CPU

histories are placed in neighboring tabs.


32

3.3.6 Flow Rates

The HPLC pumps used for the solvent are equipped with a digital display for

flow rate read-out but the Milroyal D pumps are not. As such the precursor

flow rates are initially unknown. The stroke volume of a pump is determined

by the current input fed into its actuator. This current level could a priori be

used directly to calculate the flow rate, but this is not feasible, as the pump

output is pressure dependent. A setting of 10% pump capacity may produce 5

mL/min at ambient pressure, while at 250 bar a setting of 20% is required to

produce the same flow rate. In addition, the pump head may trap small

pockets of air or sediments, which would temporarily reduce the stroke

volume. From experience equal flow rates usually never translates into equal

current levels when multiple Milroyal D pumps are used.

Flow rates are a common problem in process technology, and are usually

determined by flowmeters. While these provide instant flow measurements,

they add extra connections, dead-volume and cost, and might eventually fail

due to corrosion or clogging as they are in direct contact with the precursors.

The solution was to place each bottle on a scale (Kern 440-47N), collect the

data from its serial port and perform linear regression over the last six data

points sampled every five seconds. Dividing mass-loss by density and

changing sign (a positive flow rate is due to a negative change in mass) gives

the flow-rate. This method is very robust for evaluating long term flow

stability although each flow rate data point is an average of the last 30

seconds.

The Kern scales are equipped with DB9 ports for RS232 communication, but

serial communication is unsuitable for data transmission over distances of

more than a few meters. The data from each scale therefore had to be

collected by a serial server (MOXA NPort 5610-8-DTL) placed in close

proximity and connected to the Local Area Network. The PLC communicates

with the serial server by knowing its IP address. All five scales are connected

to the server.

3.3 Construction

33

3.3.7 LabVIEW

The software behind the GUI and the software running on the PLC are both

applications developed using LabVIEW. This is a graphical programming

language in which data flows are generated by wiring inputs and user-defined

variables to functions and logical operators. This kind of visual command

representation is advantageous compared to text-based programming, as it

provides an overview of the structure, and the data flows are fairly intuitive.

Additionally, it eliminates problems with syntax errors such as missing

brackets which enables easy debugging. A course in data acquisition was

taken at AU to learn the basics of LabVIEW.

A simplified example of a part of the source code running on the PLC is

shown in Figure 3.8. The outer border represents a while-loop which will

execute the loop contents every 500 milliseconds until the Stop Loop Boolean

changes from false to true. A Loop Count indicator counts the number of loops

executed which can be used to monitor whether a process is running or has

stopped. The loop content starts with the two inputs, AI9 and AI10. These

represent analogue current input positions in the PLC hardware to which the

two pressure sensors have been connected. The current input is automatically

converted from mA into bar. The two pressure inputs are wired to a

comparison function which returns the maximum value. This value is then

compared with 325: If the pressure is higher than 325 bar the DO0 Boolean

will change status. The DO0 is a digital output in the PLC. It is connected to

an electronic switch in the power circuit for the pumps and heaters in Figure

3.6. The switch is normally open and only closed if input<325 is true. This

enables fast auto-shutdown of the apparatus in case of uncontrolled pressure

buildup which protects the user and apparatus from each other.

It may have occurred to the reader that operation of the pumps could also

have been facilitated from the software by generation of analogue current

signals from the CRIO into each pump actuator instead of using the black

control box with its manual turn dials. This was a deliberate choice as

operation should be aided by but not dependent on digital control. Although


34

the PLC must be switched on for the heaters and pumps to be on a closed

circuit the GUI is, however, not a prerequisite for running an experiment.

Figure 3.8 | Simplified block diagram (source code) for part of the LABVIEW application

running on the PLC. Two pressure inputs (AI9, AI10) are read, the maximum value is

found which is compared to 325 bar. A digital output DO0 changes state if the pressure is

too high.

3.3.8 Safety

The elevated pressures and temperatures used in hydrothermal flow

synthesis pose a risk that must be addressed and reduced for a safe working

environment. The hardware is rated by the manufacturer to conditions more

severe than the operating conditions. The increased wall-thickness of MP

tubing makes it suitable for pressures up to 1500 bar, but increased

temperatures lowers the maximum working pressure. At 537°C the working

pressure rating of SS316 and Inconel-625 MP tubing drops to 1000 bar. For

Hastelloy the maximum working pressure for 1/8’ tubing at 300°C is 480

bar.56 Besides making the apparatus robust enough for the harsh synthesis

conditions a total of three safety features are implemented in the reactor

system:

3.4 Sub-conclusion

35

- A burst diaphragm is installed in a cold section of the apparatus which

is standard in high pressure setups. It consists of a thin rupture disc in

contact with the pressurized fluid and ambient air on either side. It

constitutes the weakest point of the apparatus with a burst rating of

515-535 bar, and in case of emergency it directs the reactor contents

away from the user.

- Pressure is generated by the pumps; at 400 bar the HPLC pump will

automatically fail due to an internal pressure sensor, and at 450 bar an

internal PRV will open in the Milroyal D pumps and spill hydraulical

oil into an outer circuit.

- Automated shut-down of pumps and heaters if pressure exceeds 325

bar. The alarm/pressure control loop is embedded in the PLC and is

updated every 0.5 seconds. The 325 bar setpoint can be increased

manually by the user if necessary. Triggered alarms generate a

notification email to super-users. Twelve months after commissioning a

total of 58 alarms had been registered.

Prior to commissioning a shake-down was performed to verify system

integrity. The pressure was slowly increased up to 400 bar, and then

decreased. The temperature was then raised up to 450°C in steps of 50°C, and

for each temperature step the pressure was increased to 400 bar and

decreased again. Leaks were observed for a few connections which were

tightened accordingly, but otherwise shakedown was successful.

3.4 Sub-conclusion

A novel dual-stage hydrothermal flow reactor was designed and successfully

constructed. The formation of core and shell materials is divided into separate

reactor zones to prevent a mixed product. The temperature profiles of these

two zones differ because the growth mechanism of core and shell are different.


36

By design the slow heating in the secondary reactor will promote

heterogeneous deposition on a core particle.

Several processes occur throughout the reactor system and automatic

monitoring provides an extensive system overview and enables the reactor to

be operated safely by a single user. Commissioning of the reactor was done

using only the primary reactor in a single-stage setup which is the topic of

Chapter 5. The full dual-stage setup is employed in Chapter 6 where the

choice of design is evaluated. Before discussing these results the

characterization techniques used to obtain them are described in the

following chapter.

37

Chapter 4

Characterization Techniques

When a synthesis product has been collected a natural first question is:

“What is it?” This chapter provides a description of the main instruments and

analysis tools used to investigate the characteristics of the collected product

and answer that question. Different tools provide different information about

the product, and by combining several characterization techniques the

researcher is able to gain a better understanding of the synthesized product.

Central to this dissertation are the combination of diffraction and microscopy:

the former to determine the crystalline phases in the collected product and

the latter to discover whether the phases form a nanocomposite or mixed

particles.

Chapter 4 Characterization Techniques

38

4.1 X-Ray Diffraction

X-rays were first discovered by Wilhelm Röntgen in 1895 as he was

experimenting with discharge tubes and noticed a faint yellow gleam from a

nearby fluorescent screen. This new radiation penetrated the paper cover of

the tube and its penetration ability was quickly demonstrated by a

photograph of his wifes hand showing only her bones as the rays passed

easily through the tissue. It was later demonstrated by von Laue and

subsequently by W.H. Bragg and W.L. Bragg that this radiation could also be

used for obtaining diffraction patterns from a crystal. These discoveries laid

the foundations for the field of crystallography which constitute a central part

of materials analysis.57

4.1.1 Generation of X-Rays

X-rays are electromagnetic radiation with wavelengths of the order of 10-10 m.

They can be generated by emission of photons produced by electronic

transitions within the atom. These transitions can be induced by

bombardment of matter with a high-energy beam of electrons. By applying a

large potential difference between a heated tungsten filament as anode and a

target as cathode the electrons will accelerate towards the target and produce

a spectrum of photons as the electrons interact with the target (Figure 4.1A).

The spectrum consists of two components. One part is a continuous

distribution of white radiation as electrons are slowed down and eventually

stopped also known as Bremsstrahlung. Superimposed on this spectrum are a

number of sharp peaks associated with electronic transitions of the target

atoms. Figure 4.1B shows how a core-electron is knocked out by an incoming

high-energy electron. This leaves a vacancy in the ionized atom, and an outer

orbital electron will quickly drop down to occupy the vacant position. This

occurs with the simultaneous emission of a photon which is equal in energy to

the energy difference of the two orbitals. Two main electronic transitions are

observed in Figure 4.1A: Kα from the 2p→1s and Kβ from the 3p→1s


39

transition. The Kα transition is the most intense, and consists of two discrete

energies Kα1 and Kα2 from the two different spin states. These photonic

emissions are unique to each chemical element. Cu is often used as target

with Kα1 = 1.54051 Å, but other targets may be chosen to obtain a different

wavelength for diffraction.58

Figure 4.1 | A) Illustration of the X-ray spectrum from electronic interactions with a

target. Inspired by figure from West.58 B) A high-energy electron ejects a 1s core electron

triggering an electronic transition during which x-rays are emitted. Inspired by figure from

West.58

4.1.2 Bragg’s Law of Diffraction

For crystalline materials the atoms are periodically arranged in a crystal

lattice. In the Bragg description of diffraction the solid is considered as

constructed of semi-transparent mirror planes (Figure 4.2), and a

monochromatic X-ray beam incident on the sample at an angle θ will be

reflected by these planes. For those angles θ in which the distance XYZ is

equal to a whole number of wavelengths λ constructive interference of wave 1

and 2 will occur and form a diffracted beam. This may be expressed as

𝑋𝑌𝑍 = 𝑛𝜆 = 2𝑑ℎ𝑘𝑙sin(𝜃ℎ𝑘𝑙)


40

Here dhkl is the distance between two crystal planes of Miller indices (h,k,l).

By noting that the nth order of diffraction of a given (h,k,l) plane is

equivalent to a first order of diffraction of (nh,nk,nl) we can exclude n. This

can be summarized as

Bragg’s law 𝜆 = 2𝑑ℎ𝑘𝑙sin(𝜃ℎ𝑘𝑙)

Reorientation of the sample will divert the diffracted beam, and for a

polycrystalline sample with crystallites of random orientation the combined

diffraction beams will form a series of diffraction cones. The diffraction

intensities can be obtained by a detector, and the plotted data is known as a

diffraction pattern, where each peak can be correlated to an inter-planar

distance in the crystal structure.58 Diffraction data was obtained using a

Rigaku SmartLab Diffractometer with a rotating anode, and data was

collected from a movable 1D detector.

Figure 4.2 | Illustration of Bragg’s law. Atoms in plane A and B scatter the incident beam.

A diffracted beam is formed if the distance XYZ is equal to a whole number of wavelengths.

Inspired by figure from West.58


41

4.1.3 Structural Analysis

Rietveld Refinement is a method used to treat the crystallographic data. It is

a least-square-method developed by Hugo Rietveld, and the algorithm is

available in the FullProf software package.59 The algorithm uses a structural

model as input with a range of variables to fit a diffraction pattern. The

function

𝑀 = ∑ 𝑤𝑖,𝑜𝑏𝑠(𝑦𝑖,𝑜𝑏𝑠 − 𝑦𝑖,𝑐𝑎𝑙𝑐)2

is minimized, where yi,obs and yi,calc are the observed and calculated intensities

and wi = 1/σi² is the weight of each yi,obs intensity at step i. The calculated

intensities for a crystalline phase are obtained by summing over all

contributing hkl Bragg-reflections

𝑦𝑖,𝑐𝑎𝑙𝑐 = 𝑦𝑖,𝑏 + 𝑆 ∑ 𝑗ℎ𝑘𝑙𝐿ℎ𝑘𝑙𝑂ℎ𝑘𝑙|𝐹ℎ𝑘𝑙|2Ω𝑖,ℎ𝑘𝑙

ℎ𝑘𝑙

where yi,b is the background intensity, S is a scale factor related to the

diffracting sample volume, jhkl is the multiplicity factor of the hkl reflection, L

includes both the Lorentz and polarization factor related to the geometry of

the experimental setup, O scales intensities in case of preferred orientation, F

is the structure factor and Ω describes the peak profile function. The

background points yi,b can be either by linear interpolation of selected

background points between the Bragg peaks or a polynomial function.

The structure factor, Fhkl, is a mathematical expression of how incident

radiation is scattered by the electric field of a material. It is obtained by

summation of the scattering by all atoms, i, in the unit cell for a given hkl-

reflection.

𝐹ℎ𝑘𝑙 = ∑ 𝑓𝑖𝑒𝑥𝑝 (−𝐵𝑖

𝑠𝑖𝑛2𝜃ℎ𝑘𝑙

𝜆2) 𝑒𝑥𝑝(2𝜋𝑖(ℎ𝑖𝑥𝑖 + 𝑘𝑖𝑦𝑖 + 𝑙𝑖𝑧𝑖))

𝑢𝑛𝑖𝑡 𝑐𝑒𝑙𝑙

𝑖=1


42

where f is the atomic scattering factor which decrease with scattering angle,

B is the Debye-Waller factor which describes the thermal vibrations, and

(xi,yi,zi) are the fractional coordinates of the atoms in the unit cell.60

PXRD is a strong bulk technique for product investigations. The Rigaku

diffractometers allows fast sample analysis and the majority of the collected

data in this dissertation is diffraction based. Its main limitation is the

inability to detect non-crystalline phases.

4.2 X-Ray Fluorescence Spectroscopy

XRF is a spectroscopic technique that allows the identification and

quantification of chemical elements present in a sample. The sample is placed

in a beam of X-rays which may excite the electrons and ionize the atoms. If

the photon energy is high enough it will eject a tightly bound core-electron,

and an outer electron will fall into the vacant orbital by emission of florescent

radiation characteristic for the irradiated atom.

The electronic transitions may only be generated if the photon energy is

larger than the orbital energy, but the number of transitions decrease with

photon energy. Moreover the absorption lines may cover a wide spectrum -

the core electrons of heavier atoms are more tightly bound, and their Kα will

be at higher energies than for lighter elements. In order to maximize the

number of observed absorption lines and increase detection sensitivity the

sample is therefore hit by X-rays of different energies using different targets.

The XRF spectrometer is a Spectro Xepos with 4 different targets.

XRF is a fast bulk technique but is only able to identify elements heavier

than sodium. Additionally, the qualitative yields suffer if a sample with a

high concentration of a single element is analyzed e.g. a lump of steel.

Similarly, low sample volumes lead to loss of radiation and the calculated

yields are not trustworthy. Fortunately, the former problem is insignificant as

most samples were metal oxides. Although oxygen and other light elements

4.3 Scanning Transmission Electron Microscopy

43

cannot be identified they will still absorb radiation and prevent detector

saturation. The latter problem is also insignificant as the primary use of XRF

in this dissertation is to find the atomic ratio of different metals. Generally,

several grams were synthesized for each flow synthesis which is plenty for

XRF analysis.

4.3 Scanning Transmission Electron Microscopy

An electron beam produced by acceleration of electrons in an electric field is

used to interact with a sample and generate an image, and magnetic

condenser lenses are used to manipulate the beam path in the same way as

glass lenses are used in light microscopes. The electron beam is scanned

across the surface, and an image is formed by relating the detected signal

with the beam position. Different detectors can be used. In bright field the

detector is placed immediately below the sample, and an image is generated

by the electrons passing through the sample. This imaging method is used in

Transmission Electron Microscopy (TEM), and the contrast is due to sample

absorption.

In Scanning Transmission Electron Microscopy (STEM) a High Angle

Annular Dark Field (HAADF) detector can be used to detect the scattered

electrons. This signal does not include the transmitted beam, and it is

therefore known as a dark field technique. It is similar to SEM but in STEM

the HAADF detector is placed beneath the sample, and only the scattered

electrons transmitted through the sample are detected. This allows for higher

spatial resolution, as the electrons are scattered and collected over a large

angle, which makes the contrast very sensitive to variations in atomic

number.

The STEM is a TALOS F220A from FEI operated at 200 kV. It is equipped

with an Energy Dispersive X-ray (EDX) detector. The X-rays emitted from

electronic transitions are used to identify and quantify the elemental


44

composition of the investigated sample. These elemental distributions can be

produced as an image.61

STEM is a highly advantageous technique for determination of whether two

phases form a hybrid material, but it suffers from two major drawbacks.

Firstly, it is not a bulk technique as only a minute fraction of the total

product is used for analysis. Secondly, it is very time consuming.

4.4 Fourier Transform Infrared Spectroscopy

FTIR is an absorption technique in which a sample is irradiated and a

spectrum of the transmitted radiation is recorded. This spectrum will differ

from the background (sample-free) spectrum due to sample absorption. The

energies of infrared radiation are comparable to the energies of molecular

vibrations and absorbed wavelengths are due to bending and stretching of

chemical bonds in the sample. The obtained spectrum is therefore used to

identify the presence of different chemical groups e.g. hydroxides, alcohols etc.

but also metal-oxygen bonds can be detected although not at great resolution.

FTIR was acquired using a Nicolet Avatar 380 from Thermo Electron

Corporation.62

The main use for FTIR in this dissertation was for determining the presence

of silica. SiO2 is non-crystalline and diffraction was therefore not a suitable

technique for bulk investigations of silica-containing samples.

4.5 Sub-conclusion

Combination of several characterization methods provides much information

of the collected product. By data comparison the results can be corroborated

to form an unambiguous conclusion. Each method has its strengths and

weaknesses and product characteristics should not be based solely on one

technique.

45

Chapter 5

Commissioning

Once the reactor was constructed and shake-down was performed to verify

the structural integrity a range of nanoparticles were synthesized as part of

the commissioning. This chapter discusses the synthesis results which formed

the basis of reactor commissioning. The purpose was to confirm that the

reactor with all its components was functional and suitable for hydrothermal

flow synthesis of inorganic nanoparticles. The single-stage setup was used for

this purpose, and the synthesis results were compared with previous works

on the old flow reactor. The results from commissioning are reported in

Publication I which also discusses the choice of design examined in Chapter 3.

Chapter 5 Commissioning

46

5.1 TiO2

TiO2 was selected as studies performed on the old flow reactor by Mi et al.63

provided a broad parameter space for direct comparison of the two reactors as

syntheses performed on different reactor systems are not guaranteed to

behave identically. Titania is a highly versatile semi-conductor material used

in a broad range of applications including batteries, photocatalysis, dye-

sensitized solar cells but also as white pigment in paints and foods.64 It has

therefore been intensely researched, its formation chemistry is well known,

and is an excellent model material to characterize the reactor performance.

Upon successful commissioning of the single-stage setup the full performance

of the dual-stage reactor was demonstrated with a proof-of-concept synthesis

of a TiO2@SnO2 nanocomposite which will be addressed in Chapter 6. The

main synthesis results from commissioning form the basis of Publication I

Titania can crystallize in several different phases among which the most

common are rutile, anatase and brookite. Although rutile (ICSD 167958) is

the thermodynamically most stable phase, phase-pure anatase (ICSD 92363)

was obtained for all of the investigated syntheses during commissioning.

Equal flow rates of 10 mL/min for solvent and precursor were used

throughout commissioning, and the temperature of solvent main heaters was

kept identical to reactor heaters.

The single-stage was also in use for production runs of Strontium hexaferrite

(SrFe12O19) performed by Mathilde Múzquiz. The purpose of these was for

magnetic applications for which a high throughput was necessary. This

required a reliable reactor performance for at least 8 hours per day which was

used to evaluate long-term reactor stability in continuous operation. After

several days of continuous operation a fuse broke in the electrical board for

one of the precursor pumps, but otherwise the production run was successful.

Other single-stage synthesis work includes AlOOH and γ-Al2O3 by Sanna

Sommer and TiO2 on graphene support for photocatalysis by Pedro

Magalhães.

5.2 Flow synthesis of TiO2

47


A solution of Titanium tetraisopropoxide (TTIP) in isopropanol (isopropyl

alcohol, IPA) was used for titania precursor, and the solvent was

water/isopropanol mixtures of different water contents. Each synthesis began

by pumping solvent from the HPLC pump and pure isopropanol from the

precursor pump. Precursor flow rates are adjusted by the controller box until

at the desired level during. Once the heaters reach a steady state the

precursor is subsequently injected by diverting the glass valve inlet from IPA

to TTIP. The time from a precursor enters the system until the product

emerges from the PRV is roughly 5 minutes and twice that for the dual-stage

reactor.

The collected product was a suspension which was centrifuged into a

sedimented layer of TiO2 and a clear supernatant above. The supernatant

was removed, and the particles were washed with water/ethanol repeatedly.

The wet slurry was placed in a vacuum oven overnight at temperatures of 50-

70ºC. The dried powder was crushed with a pestle in a mortar. This procedure

was the same for all collected products.

The influence of reaction temperatures, water content in the solvent, and Ti

precursor concentration on crystallite size and crystallinity was investigated.

An increase in water contents is seen to decrease the crystal size in Figure 5.1.

For 100% water the crystal size is approximately 6 nm which was also

observed by Zhang et al.65 High water content means that more water is

available for hydrolysis of TTIP which means more crystals are formed and

the average crystal size therefore decreases. The crystal sizes for the different

TTIP concentrations are all found to converge at 6 nm at high water contents.

For lower water contents the crystal size varies from 12 to 15 nm. As water is

no longer available in excess the nucleation step becomes increasingly

sensitive to fluctuations in the local water concentration in the mixing piece.

The dual pump head HPLC provides a non-pulsating solvent flow whereas

the Milroyal D pumps have a single pump head. The local water

concentration in the mixing piece therefore decreases with each precursor


48

injection and subsequently increase as the pump is refilled. Large variations

in crystal size at low water contents were also observed by Mi et al.63 who

observed crystal sizes of 13-18 nm in this region. Similar crystal sizes were

obtained by Kawasaki et al.42

Figure 5.1 | The average crystallite size decreases with water content for all

concentrations.

The nanoparticles synthesized on the old flow reactor generally have a

slightly larger crystal size than those synthesized on the new reactor. This is

likely due to the heavily pulsating pneumatic pumps which decrease the local

water content when a precursor pulse is pumped towards the mixing point.

However, the calculation methods differ as Mi et al.63 used the Scherrer

equation based on a single diffraction peak fitted with a Lorentzian function

while Rietveld refinements were used in the current experiments.

Different TTIP concentrations were found to have no influence on the crystal

size in the investigated concentration range of 0.125 M to 1.0 M at a

temperature of 350°C as seen in Figure 5.2. At a maximum concentration of

1.0 M TTIP the viscosity greatly increased, and the product was a paste. This


49

resulted in large fluctuations in the pressure, and any differences in product

characteristics are likely due to the instable pressure and flow. The reactor is

more prone to clogging when the concentration is increased. This was

observed at one point when a pump stalled, and the synthesis had to

temporarily halted. When the synthesis was resumed the stagnant titania

paste had sedimented onto the hot reactor walls. A drill was used to clear the

reactor tube, but the walls could not be cleared of remaining sediments. The

tube was eventually discarded and replaced.

Figure 5.2 | No change in crystal size is observed for the different concentrations.

Figure 5.3 and Figure 5.4 show how the crystal size and crystallinity,

respectively, increase with temperature. For these syntheses the

concentration was kept at 0.25 M.66 This was also observed in the previous

study. While the trend of increasing crystallinity with temperature is in

agreement with previous results absolute values should not be compared. In

that study CaF2 was used for an internal standard. The salt is very

hygroscopic and absorbed moisture would reduce the crystallinity of the


50

standard and increase the calculated TiO2 crystallinities. In the current study

silicon was used as a standard to avoid this.

Figure 5.3 | The average crystal sizes increase with temperature up until a maximum of

approximately 15 nm.

An interesting feature is that 15% higher crystallinities are observed for 50%

water than for 100% water and this shift is independent of temperature. The

heat capacity of water is higher compared to isopropanol which means a

water-rich reactor medium will take longer to heat up. Another reason could

be due to the arrangement of the heaters; the tubing connecting the second

solvent heater to the mixing point is insulated, but a significant heat

exchange with the ambient air is unavoidable, and this lowers the

temperature of the solvent stream just prior to mixing. The critical

temperature of isopropanol is 237°C and a 50/50 water/isopropanol medium

will therefore experience a larger volume expansion at hydrothermal

temperatures compared to pure water. The volume expansion reduces the

density and causes the flow speed to increase. A faster flowing fluid will have

less time to be cooled, and this reduces the drop in solvent temperature. The

5.3 Sub-conclusion

51

50% water stream is therefore hotter than the 100% water stream upon

mixing which increases the crystallinity of the 50% water samples.

Figure 5.4 | Higher reaction temperatures increase the crystallinity. Higher crystallinities

are observed for lower water content.

Conversely, the residence time in the reactor for the 50% water samples is

decreased which should lower the crystallinity. This appears, however, to be

negligible, and most of the formation chemistry and crystallization takes

place in the top of the reactor. Although absolute crystallinities in the two

studies are not directly comparable the values are very similar despite the

fact that the reactor tube of the old flow reactor is 0.5 m while the dual-stage

employs 1.1 m reactor tubes.

5.3 Sub-conclusion

Titania was shown to crystallize in the anatase crystal structure with an

average crystallite size of 5-15 nm depending on the selected synthesis


52

parameters. Increased reaction temperature was shown to increase the

crystal size and the crystallinity of the samples, but up to a maximum of 15

nm. Increased water content in the solvent stream promotes nucleation which

decreases the crystal size, whereas the changes in concentration had no

notable effect on the synthesis parameters. This is advantageous as it allows

a high throughput of up to 50 g/hour without changing the product

characteristics.

With the reproduction of powders with product characteristics in excellent

agreement with previous works the commissioning of the reactor was a

success. Next step is to switch the configuration from single-stage to dual-

stage setup to investigate if the dual-stage design can promote the formation

of nanocomposites under hydrothermal conditions. This is the topic for the

next chapter.

53

Chapter 6

Formation of a TiO2@SnO2

Nanocomposite

Hydrothermal synthesis of phase pure anatase using the single-stage reactor

provided a successful commissioning of the reactor as described in the

previous chapter. Following this the dual-stage setup was employed for the

proof-of-concept synthesis of a nanocomposite. This chapter discusses the key

synthesis parameters involved in heterogeneous deposition of SnO2 on TiO2.

The main results from this experiment are reported in Publication II.

Synthesis of SnO2 in the presence and absence of TiO2 is a fast way of

determining if TiO2 induces any changes in SnO2 growth. It is shown how a

TiO2@SnO2 nanocomposite can be synthesized at low pH, but that at high pH

a mixed product is formed. Ball-milling of premade powders did not produce a

hybrid, which confirms that composite formation is due to growth of the tin

precursor directly on the anatase nanoparticles and not by simple

aggregation of the two particles.

Heterogeneous deposition was shown to be dependent on tin precursor. The

precursor concentration may be increased for a better coating but higher

concentrations will lead to agglomeration.

Chapter 6 Formation of a TiO2@SnO2 nano composite

54

The dual-stage capability for nanocomposites was demonstrated by removing

the PRV from the single-stage flow reactor and connecting the primary

reactor to the secondary reactor for formation of a TiO2@SnO2 nanocomposite.

The anatase suspension synthesized in the primary reactor was cooled down

and mixed with a 0.1 M SnCl4 stream, and the contents led into the secondary

reactor for coating with tin dioxide.

6.1 Temperature and pH

By division of primary and secondary material formation into separate

reactor environments the dual-stage allows the pH of the medium to be

changed. This feature was demonstrated with the introduction of NaOH by a

third precursor pump. Upon hydrolysis the pH drops due to formation of

hydrochloric acid

𝑆𝑛𝐶𝑙4 + 4𝐻2𝑂 → 𝑆𝑛(𝑂𝐻)4 + 4𝐻𝐶𝑙

To preserve flow rates the third precursor pump was used to pump water for

the syntheses performed at low pH. All three precursors (Titanium

tetraisopropoxide (TTIP), SnCl4 and NaOH or H2O) were pumped at 4

mL/min, and hot solvent mixing was used in both reactor stages with water

pumped at 10 mL/min. The chloride concentration in the secondary reactor

was then 0.05 M and the pH either 1.3 or 13.7 depending on whether H2O or

NaOH was used as third precursor. Hot mixing in the secondary reactor was

used to test the system at maximum capacity and to gain insights on the

significance of hot mixing versus gradual heating without hot mixing. The

configuration shown in Figure 3.1 was used.

Mixing of the tin chloride with base formed a viscous gel which could not be

pumped by the pump, and the two precursors had to be added separately.

Due to the increase in viscosity the SnCl4 solution was mixed with the

anatase suspension first, while the base was added last. This caused the pH


55

first to decrease and then increase which is not ideal as it could cause

agglomeration of the titania nanoparticles. Nevertheless, this procedure was

chosen as the lesser evil compared to rapid gel formation upon introduction of

tin chloride into an alkaline titania suspension which would hinder mixing of

Sn4+ with the titania nanoparticles.

The anatase reaction temperature in the primary reactor was kept fixed at

350°C, while the temperature in the secondary reactor and secondary solvent

heaters were increased from 150°C to 300°C. This was performed for the

acidic and the alkaline series i.e. using water or NaOH as third precursor. For

each nanocomposite synthesis a titania-free reference was also synthesized.

These references are denoted Ø&SnO2 to emphasize that tin dioxide is formed

in the secondary reactor. Conversely, SnO2 formed in the presence of anatase

nanoparticle is labelled TiO2&SnO2. If the two phases forms a nanocomposite

then that material is denoted TiO2@SnO2.

For all experiments SnO2 formed phase pure cassiterite which is rutile-

structured. Figure 6.1 shows PXRD data for the SnO2 references at acidic

conditions. Rietveld refinements reveal that between 150°C and 200°C the

average SnO2 crystallite size changes from 1.4 nm to 1.6 nm. However, for

temperatures at 250°C and 300°C the SnO2 peak shape sharpens significantly

but with extended tails. The peak profile was treated as a superposition of a

narrow and a broad component, indicating the presence of both large and

small nanocrystals i.e. a bimodal size distribution. To describe this a second

phase was included in the refinements to account for these larger crystals.

Rietveld analysis revealed that SnO2 crystals of 48 and 35 nm appeared

alongside smaller 2 nm nanocrystals at 250°C and 300°C. At 300°C near half

of the crystalline fraction consisted of these larger crystals.


56

Figure 6.1 | PXRD of Ø&SnO2 at 150°C to 300°C at low pH. The

peak shape sharpens with temperature indicating both large and

small crystals are present in the product.

The bi-modal size distribution is largely a consequence of the heating profile

employed in the secondary reactor. The secondary hot mixing forms some

nucleation sites, but the flow rate is too low and heating is too slow for the

entire precursor to be consumed in this step. These SnO2 nuclei then enter

the reactor zone and continue to grow into large extended crystals in the

presence of the high concentration of unreacted tin precursor. Further down

the reactor the temperature is eventually high enough for homogenous

nucleation to occur, and this leads to formation of the smaller 2 nm

nanoparticles. These particles are formed further down, at a lower tin

concentration and shorter residence time which limit their final crystal size.

This presence of both small and large nanoparticles was confirmed by TEM

(Figure 6.2).


57

Figure 6.2 | TEM of Ø&SnO2 at 300°C at low pH.

The product consists of small isotropic nanoparticles of

a few nanometers and large extended rods.

Ostwald ripening has previously been suggested from in-situ studies as a

growth mechanism for the formation of large SnO2 crystals at temperatures

above 250°C.67 This is estimated to be a negligible effect in the dual-stage

reactor due to the short residence time and low concentration of the tin

precursor. A growth mechanism of oriented attachment of smaller SnO2

nanoparticles into larger nanorods was suggested by Firooz et al.68 which is in

good agreement with the observed morphology in this study.

In the presence of titania the SnO2 crystal size remain at 1.4 to 1.7 nm from

150° to 200°C. At higher temperatures only 2 nm crystals are present and

large crystals are no longer formed. At 250°C a small sharp peak is present in

the product, but this disappears at 300°C and was likely just residual powder

from the reference synthesis still present inside the reactor tube leaching into

the product stream. From PXRD (Figure 6.3) it is evident that TiO2 stabilizes

the SnO2 crystallites from further growth. This was confirmed by STEM-EDX

(Figure 6.4) which showed SnO2 densely deposited on TiO2 nanoparticles.


58

Figure 6.3 | PXRD of TiO2&SnO2 at 150°C to 300°C at low pH. No

sharpening of the SnO2 peaks is observed.

Figure 6.4 | A) STEM of a TiO2@SnO2 nanocomposite at low pH and 300°C. SnO2

decorate the surface of the TiO2 particles. B-E) STEM-EDX of highlighted region. Shell

consists of SnO2 particles in agreement with crystallite size.

The evolution in crystal size is shown in Figure 6.5, and the calculated crystal

sizes are in good agreement with the particle sizes observed from electron

microscopy. The anatase crystal size remains fixed at 8-9 nm throughout the


59

experiments which indicates that it is not affected by the deposition reactions

in the secondary reactor. The formation of extended SnO2 nanorods means

that the assumption of isotropic particles made in Rietveld refinements is no

longer valid, and the lengths of the nanorods are also seen to differ

significantly. The calculated average crystal size of these larger particles is

therefore associated with a significant error. However, the purpose is not to

provide an accurate description of the SnO2 references in terms of size

distribution and aspect ratio but rather to show that a significant change

occurs in the SnO2 peak profile and size distribution when TiO2 is present.

Figure 6.5 | Crystal sizes of TiO2 and SnO2. Larger SnO2 crystals emerge at 250°C and

above when TiO2 is no longer present in the reaction mixture.

Substitution of Sn for Ti in the anatase unit cell leads to expansion of the unit

cell. From Rietveld analysis it is evident that anatase unit cell size is not

significantly affected by Sn substitution which is likely due to the short

residence time. Up to 15% Sn can be substituted into the anatase unit cell. At

higher Sn contents the rutile structure is the more thermodynamically stable

phase. Unit cell sizes of selected samples are listed in Appendix 2.


60

For the alkaline syntheses SnO2 did not form at 150°C but only at 200°C and

above. The crystal size was approximately 2 nm but the diffraction profile

was highly asymmetric towards lower angles. At higher temperatures the

profile sharpened significantly indicating the presence of large particles with

increased temperature, but the significant asymmetry made the pattern

difficult to refine. By visual comparison the SnO2 diffraction profile did not

change significantly with the presence of absence of titania. The

crystallographic data therefore suggests that SnO2 is not affected by the TiO2

nanoparticles, which implies that the product is not a hybrid material. PXRD

data of alkaline Ø&SnO2 and TiO2&SnO2 samples synthesized at 300°C are

displayed in Figure 6.6.

Figure 6.6 | PXRD of Ø&SnO2 and TiO2&SnO2 at high pH. SnO2 peaks are highly

asymmetric, and the peak shape suggests a broad size distribution.

This was confirmed by STEM (Figure 6.7) which showed the two phases as

interspersed particles with clear particle boundaries. The majority of the

anatase particles are not in contact with SnO2 indicating no preferential

interactions between SnO2 and TiO2 at high pH. The product consists of

smaller isotropic particles, extended nanorods and platelets (Figure 6.8). SnO2

has been shown to form a diverse morphology at high pH including

nanoflowers, nanorods and nanospheres.69-71


61

Figure 6.7 | STEM-EDX of TiO2&SnO2 at high pH. TiO2 particles are uncoated by SnO2

and the two phases are dispersed as a mixed product.

Figure 6.8 | STEM of TiO2 and SnO2 at alkaline synthesis conditions. SnO2 crystallizes as

isotropic particles, nanorods and nanosheets.


62

The high degree of asymmetry at lower angles indicates an extended unit cell

for parts of the diffracting volume. This may be attributed to inclusion of OH

defects, or surface relaxation from hydroxylation. At high pH the

hydroxylated surfaces of the SnO2 nanoparticles are negatively charged which

increases the bond lengths.

6.2 Composite formation

Why this difference in nanocomposite formation vs mixed particles at low and

high pH, respectively? The nucleation mechanism and kinetics of SnO2

formation are sure to be affected by the pH, but what effect should these

changes have on the preference towards heterogeneous deposition onto

titania? And why does the secondary material not form a closed shell?

At low pH Jensen et al.67 showed that SnO2 is formed from an hexaaquatin(IV)

complex but a large fraction of the tin precursor was left unreacted as tin(IV)

chloride complexes. Incomplete tin conversions have also been observed by

Fang et al.72 However, in the dual-stage reactor the tin precursor is fully

converted at low pH despite the short residence time as determined by XRF.

This is probably due to a strong thermodynamic driving force towards

heterogeneous nucleation at low pH. Composite formation is therefore most

likely due to the interactions between the positively charged anatase surface

and the charge distribution of the octahedral tin(IV) complexes that exists at

low pH.

At high pH only 20% of the tin(IV) tetrachloride precursor is converted into

SnO2 at 300°C as determined by XRF. ICP-OES measurements of the

supernatant show little difference in tin concentration of the alkaline

TiO2&SnO2 and its SnO2 reference, which supports that TiO2 plays no role in

SnO2 formation at high pH (Appendix 3). Additionally, the concentration of

the hexaaquatin(IV) complexes from which SnO2 is formed is likely reduced

at high pH due to formation of the thermodynamically more stable

[Sn(OH)6]2- species which could further reduce the SnO2 yield. These species

6.2 Composite formation

63

are negatively charged and would repel the negatively charged anatase

particles.

To optimize the conditions for composite formation the solvent pH should

therefore be chosen carefully so the secondary precursor and core particle

attract or at least do not repel each other electrostatically. A predictive tool

could be based on the concept of isoelectric point (IEP). A particle placed in a

polar liquid will become charged if one type of ion is preferentially adsorbed

on its surface. This can be by Coulombic attractions or by more specific

chemical affinities such as surfactant interactions. Highly charged particles

will repel each other and form a stable dispersion whereas particles will

flocculate if the particles have a zero net charge.73 This is the same process

that happens when milk turns sour as the slow release of acid eventually

neutralizes the negatively charged proteins. At low pH particles will be

positively charged while they will be negatively charged at high pH. The pH

where the particle displays a zero net surface charge is the isoelectric point,

and this value differs from compound to compound.

Use of this concept for composite formation should, however, be treated with

great caution as it relies on several unjustified assumptions. Firstly, the

isoelectric point of a solid oxide is probably not the same as for its ionic

precursor species or even in the nucleic state. Secondly, the isoelectric point is

affected by the adsorption of ions and molecules. Thirdly, nanoparticles are

often composed of an inner crystalline core and an outer amorphous layer

which is a problem as the IEP varies with both particle size and crystal

structure. Finally, the isoelectric point has been shown to decrease with

temperatures, and most IEPs are only determined at room temperature.74, 75

Despite this multitude of assumptions the IEP can still be a useful tool for an

a priori estimate of which compounds might prove compatible, and at roughly

what pH the synthesis should initially be performed. The usefulness of this

concept was validated in later studies although none of the results should be

interpreted as definite proof. An IEP-based choice of pH is an educated guess

and should be either advantageous or insignificant, but not disadvantageous


64

towards composite formation. The measured isoelectric points of selected

oxides are illustrated in Figure 6.9, and the values are obtained from the

compilation by Kosmulski76. Note that the values will shift to the left by up to

1 at hydrothermal temperatures. If strong electrostatic attractions are

present then it is likely that the composite morphology will be a core-shell

whereas if no electrostatic interactions are present then deposition occurs by

heterogeneous nucleation as favored by thermodynamics. This relationship

between pH, IEP and product morphology is discussed in Chapter 8 where

the factors affecting uniform core-shell or decorated core are discussed.

Figure 6.9 | Illustration of Isoelectric Points (IEP) for different metal oxides. Box width

represents the range of reported IEPs. The surface charge of a suspended particle is neutral

when pH is equal to the particle IEP. BaTiO3 is abbreviated BTO.

6.3 Concentration Effects

SnO2 was found to form a nanocomposite at low pH and high temperatures,

but the morphology was not a true core-shell structure. If the shell is

comprised of several smaller nanoparticles the shell thickness must be at

least the diameter of a single SnO2 crystal. Due to the high surface/volume

ratio of the small anatase nanoparticles the SnCl4 concentration needs to be

6.3 Concentration Effects

65

at least twice that of the TTIP precursor. This is under the assumption that

all anatase particles are 9 nm in diameter and do not agglomerate. The molar

volumes of the two oxides are similar. To provide a surplus of tin the TTIP

was therefore reduced from 0.1 M to 0.01 M as an increase in SnCl4 would

otherwise change the pH of the medium. The theoretical SnO2 shell thickness

is then 5.5 nm. PXRD data are displayed in Appendix 4, and refinements

show an increase in SnO2 crystal size to 2.8 nm. If the structure is core-shell

the shell layer should consist of about two layers of SnO2 crystals. No sharp

SnO2 peaks are present indicating that no nanorods are formed. Anatase

could not be accurately refined due to the lower content and scattering power

compared to SnO2. Only the anatase scale factor was therefore refined

STEM reveals the presence of smaller clusters of titania inside larger

polycrystalline SnO2 agglomerates (Figure 6.10). While titania is most likely

extensively coated with SnO2 the resulting agglomeration makes accurate

Figure 6.10 | A) STEM of TiO2@SnO2 with a Sn:Ti molar ratio of 10. B-D) EDX maps

shows titania engulfed by a large lump of SnO2.


66

EDX mapping of individual nanocomposites difficult as the detected signal is

a superposition of all underlying particle layers. Analysis of the few dispersed

particles in the background would be unrepresentative, as most of the product

is present as agglomerates. Agglomeration is due to a surplus of SnO2, and

the critical Sn/Ti concentration ratio that determines when this occurs is

likely related to the Sn content necessary to form a monolayer of SnO2

crystallites around the TiO2 particles.

6.4 Evaluation of Secondary Hot Mixing

Formation of nucleation sites at the top of the secondary reactor was shown to

grow large SnO2 crystals in the absence of titania. Hot mixing therefore

allows composite formation to be detectable with diffraction techniques which

is beneficial, as STEM is more time consuming and is only a local probe.

Although hot mixing may provide a strong crystallographic proof for

heterogeneous deposition the composite samples may contain ungrafted SnO2

particles as parts of the TiO2/SnCl4 mixture is heated too fast. In addition, the

increased flow rate reduces the residence time which compromises the cooler

performance at high temperatures which limits the maximum reaction

temperature. Secondary hot mixing was therefore abandoned for all

subsequent syntheses with the assumption that diffraction would still prove

able to distinguish between mixed powders and composites. This decision was

backed by subsequent synthesis work.

Each 1.1 m reactor zone is constructed of two 0.5 m separately controlled

heaters and with the removal of the secondary hot mixing the reaction

mixture proved more difficult to heat up by the top heater. At 300°C the

temperature of the top aluminium heater would slowly decrease due to the

enhanced heat exchange by the fast-flowing mixture. It was eventually

decided that if a reaction temperature of e.g. 350°C was needed then the top

heater would be set at 250°C, while only the lower heater would be set at

350°C.

6.5 Mechano-chemical processing

67

6.5 Mechano-chemical processing

TiO2 and SnO2 powders were mixed and ball-milled to investigate if

composite formation is possible once the secondary phase is fully formed. This

was to confirm that the composites are indeed grown inside the dual-stage

reactor and not simply formed by adsorption of SnO2 nanoparticles onto TiO2

surfaces after formation of the secondary phase. STEM-EDX analysis (Figure

6.11) shows a sharp surface boundary between SnO2 and TiO2. The two

phases are well dispersed, but the SnO2 are not tightly bound to the TiO2

nanoparticles as previously observed in Figure 6.4. The product is a simple

mixture of homogenous particles. A composite structure is not even formed

under the extreme milling conditions, which reaffirms that composite

formation occurs directly from the tin precursor.

Figure 6.11 | A) STEM of ball-milled TiO2 and SnO2 powders. B-D) EDX shows particles

are well-dispersed but does not form single-particle composites.


68

6.6 Sn-Precursor

Due to gas formation from degradation of isopropanol at low pH the acidic

syntheses were stopped at 300°C. The reactor was spewing hot hydrochloric

acid at elevated flow speeds for which the cooler performance was inadequate.

Additionally, the yellow-green colour of the product indicated dissolution of

the stainless steel reactor tube, and the acidic temperature series was

increased no further. Further investigation of the influence of Sn/Ti ratio on

nanocomposite shell morphology combined with the possible introduction of

surfactants to prevent SnO2 agglomeration at high Sn concentrations would

be very interesting, but due to the detrimental effect on the reactor it was

evident that chlorides were not a viable option as counter-ion. Tin(IV) acetate

was tried for a single synthesis at low concentrations, but the results were

discouraging. Moreover that precursor is 360 times more expensive than the

commonly used SnCl4x5H2O (based on prices from Sigma-Aldrich).

Nitrate salts are commonly used in hydrothermal synthesis. They are very

benign in terms of reactor corrosion and are often both soluble and cheap.

However, tin nitrate is unstable towards hydrolysis and is therefore not

commercially available. To overcome this problem a method for preparation of

a tin(IV) nitrate solution was therefore developed in collaboration with Aref

Mamakhel. Tin granules were dissolved in nitric acid with added sodium

nitrate or ammonium nitrate to catalyze the dissolution and delay SnO2

formation. Upon complete dissolution the precursor was diluted with water to

the desired molarity. The tin nitrate precursor solution was stable for several

hours but would hydrolyze overnight leaving a milky white SnO2 dispersion.

The TiO2&SnO2 experiments were therefore repeated but with a freshly

prepared nitrate precursor. Crystal sizes obtained from Rietveld analysis are

summarized in Figure 6.12A. The peak profile of SnO2 was found not to change

with the presence of anatase nanoparticle as had otherwise been observed for

the chlorides at high temperatures (although hot mixing was no longer used).

This indicated that titania did not interact with the growth of SnO2, which

6.6 Sn-Precursor

69

was confirmed by electron microscopy. STEM-EDX (Figure 6.13) showed a

mixed product with sharp uncoated titania boundaries.

Figure 6.12 | Crystal sizes as a function of temperature of TiO2 and SnO2 using different

precursors. A) TTIP and Sn(NO3)4. B) TALD and Sn(NO3)4.

Figure 6.13 | A) STEM-EDX of TiO2 and SnO2 from TTIP and Sn(NO3)4 at 350°C. No SnO2

is observed on the titania surfaces. The two phases do not form a hybrid, but a mixed

product.


70

Why this discrepancy between the use of chlorides and nitrates? At the time

this was attributed to an unstable pH which had been observed during

synthesis. Both TiO2 and SnO2 are white powders and the nanocomposite is

also white. As such the presence of both oxides cannot be easily “observed”

from the product stream. The composite synthesis-procedure was therefore

first to look for the emergence of the suspension from the uncoated core

product, and only then was the secondary precursor stream introduced. The

composite product was collected only when the exit stream was both white

and acidic as measured with pH indicator paper. At 100°C no SnO2 was

formed. At 150°C the exit stream was both white and acidic as expected, but

at temperatures of 200°C and above the white SnO2 stream turned pH

neutral. This was due to the surplus of isopropanol acting as a reducing agent

towards the nitric acid. While ethanol is known to react violently with nitric

acid, isopropanol is a secondary alcohol and is less easily oxidized. A higher

reaction temperature of >150°C is therefore necessary. The synthesis at

150°C did not yield a hybrid.

The problem of pH instability was circumvented by the use of Titanium

bis(Ammonium Lactato) Dihydroxide (TALD) which is an aqueous Ti

precursor. The use of TALD was found to generally decrease the anatase

crystal size by a few nanometers compared to the TTIP precursor as was

shown in Figure 6.12. The SnO2 calculated crystal size increased from 2 nm at

200°C up to 4 nm at 350°C independently of the presence of titania. STEM

analysis (not shown) showed a mixed product similar to Figure 6.13.

The tin granules are dissolved in surplus of nitric acid, and the nitrate

precursor is therefore much more acidic than the chloride precursor. Sodium

Acetate was added to neutralize the tin nitrate precursor for a final product

pH of 3.5. NaAc (0.9 M) was added to the primary TALD precursor as gel-

formation of the secondary tin-precursor would otherwise clog the pump. The

product was a mixed product and did not produce a composite, which could

also be attributed to the high ionic content and surfactant effects of acetate

(molar ratio of Ac:Ti is 9). It should be mentioned that no composites (TiO2@γ-

6.7 Sub-conclusion

71

Fe2O3 and TiO2@ZnO, not included in this dissertation) were ever obtained

when anatase core-particles where synthesized from TALD. This may be due

to lactic acid and its degradation products that coat the particle surfaces.

Such organic coatings can be detected by hyperpolarized solid-state NMR

which is the subject of Publication V.

The use of Sn(NO3)4 and SnCl4 illustrates the importance of the precursor

used. A better understanding of the growth mechanism of tin nitrate into the

oxide may aid in explaining the different preferences towards deposition on

anatase.

6.7 Sub-conclusion

A TiO2@SnO2 nanocomposite was successfully synthesized from TTIP and

SnCl4. Composite formation was shown to occur only at low pH, whereas

alkaline conditions lead to a mixed product and low conversion of the tin

precursor. The nanocomposite structure was confirmed by STEM-EDX, but

PXRD was also shown as a powerful tool for demonstrating secondary

nucleation by comparison of SnO2 crystal sizes for products with and without

TiO2.

Nanocomposite formation demonstrated the dual-stage capability and

justified the choice of design. Mixed phases at the interface are a common

feature for nanocomposites synthesized hydrothermally in batch, but this was

not observed due to the short residence of less than a minute. The results

demonstrated that pH can be changed between the two reactor zones if the

primary and secondary materials require different synthesis pH.

A stable tin(IV) nitrate precursor for SnO2 synthesis was developed which is

more benign towards the reactor. No differences was observed in SnO2

product characteristics when using Sn(NO3)4 or SnCl4, but only the chloride

yielded a composite with TiO2. An aqueous Ti-precursor, TALD, produced


72

smaller crystal sized than the TTIP but did not form a composite with

Sn(NO3)4.

The SnO2 decorated the TiO2 core but did not form a closed shell. In the next

chapter the synthesis conditions responsible for the core-shell morphology is

explored using silica as shell material.

73

Chapter 7

Silica

SiO2 is a classically used coating material. It does not require hydrothermal

conditions to form and composites of SiO2 constitute a large body of core-shell

publications. SiO2 nanoshells were prepared to compare the performance of

the reactor with these compounds. The results of this chapter are the basis of

Publication III.

Three different materials were used as core: γ-Fe2O, TiO2 and α-Fe2O3. From

the IEP concept introduced in Chapter 6 a non-alkaline reaction environment

should promote composite formation. However, silica growth is accelerated at

low pH. Core-shell morphologies were therefore obtained for γ-Fe2O and TiO2,

because these were synthesized near neutral pH. By contrast, α-Fe2O3 was

synthesized at low pH and gave a non-uniform silica layer.

The results show that core-shell morphologies can be induced by controlling

the pH of the primary suspension. If core-shell formation is a simple matter of

pH control then they become much easier to synthesize hydrothermally.

Predictions on how to select the correct synthesis pH should then be based on

IEPs which is the subject of Chapter 8.

Chapter 7 Silica

74

With the synthesis of TiO2@SnO2 nanocomposites a wide variety of precursors

was screened for nanocomposite formation. To limit the number of unknowns

titania was used extensively as core-material because of its well-known

chemistry and predictable behavior in hydrothermal flow reactors. The choice

of core material was later expanded to also include iron oxides which are also

well understood materials and highly interesting for applications. The

current chapter discusses these different core materials as templates for silica

deposition

7.1 SiO2 Shells

Silica has been used for the coating of several core nanoparticles including

TiO2,77-80 Au,81, 82 Fe2O3,83, 84 Fe3O485-88 and ZnO89, 90. Silica deposition is

known to protect the core, improve the dispersion stability, and it allows for

further functionalization with organic molecules for use as a biosensor or in

medicine.91 Deposition of silica on a photocatalytic material reduces the

performance of the core with increasing shell thickness, but Feng et al.92

demonstrated that for thin silica shells (~5.5 nm) this effect is negligible, and

the composite retains most of its photocatalytic performance. Control of the

shell thickness is therefore of great importance.

SiO2 can be readily formed by hydrolysis of the tetraethyl ortosilicate (TEOS)

followed by condensation reactions in a typical sol-gel process.93 The synthesis

can be performed at ambient conditions and due to the low scattering power

compared to most transition metal oxides the composites are easily observed

in TEM. Typical silica shells obtained from batch experiments are in the

range of 20-100 nm, but shell thickness of approximately 4 nm has previously

been demonstrated.80

Three nanocomposites were synthesized: γ-Fe2O3@SiO2, TiO2@SiO2 and α-

Fe2O3. γ-Fe2O3 was synthesized from 0.1 M Ammonium Iron Citrate (AIC),

TiO2 was synthesized from 0.05 M TTIP and α-Fe2O3 from 0.1 M Fe(NO3)3. A

lower TTIP concentration was used because the anatase particles are smaller

7.1 SiO2 Shells

75

than the iron oxides and will generate a larger specific surface area. The

secondary precursor was 0.2 M TEOS in isopropanol as it hydrolyses slowly in

water. The iron oxides were formed from water pumped at 15 mL/min and an

iron precursor flow rate of 5 mL/min. For TiO2 the solvent was a 1 mM acetic

acid solution pumped at 12 mL/min in order to lower the pH to approximately

4. All three core particles were synthesized at 350°C in the main reactor

while the secondary reactor was kept at 100°C. The dual-stage was

reconfigured using only three pumps as illustrated in Figure 7.1. No secondary

hot mixing was used and with only two reactants the residence time in the

secondary reactor is maximized. This experimental setup was the most widely

used, and is referred to in the following as the “standard” dual-stage setup.

Figure 7.1 | Standard dual-stage configuration used for silica shells and other

nanocomposites. The secondary reactor content is heated gradually without the use of a hot

mixing in this setup.

Chapter 7 Silica

76

PXRD data of all three nanocomposites are shown in Figure 7.2 along with the

uncoated sample.

Figure 7.2 | PXRD and Rietveld refinements of silica coated and uncoated metal oxides. A)

γ-Fe2O3@SiO2 and γ-Fe2O3 B) TiO2@SiO2 and TiO2 C) α-Fe2O3@SiO2 and α-Fe2O3.

7.1.1 γ-Fe2O3@SiO2

Figure 7.2A shows the diffraction patterns using the AIC precursor and the

average crystal size is 18.8 nm and 18.6 nm for the uncoated and the coated

sample respectively. These crystal sizes are in good agreement with results

obtained from in-situ PXRD studies.94 Maghemite (γ-Fe2O3, ICSD 247036)

and magnetite (Fe3O4) are both spinel structures and crystallographically

very similar, but the AIC precursor has been shown to produce mainly

maghemite.95

TEM of the phase-pure maghemite nanoparticles prior to coating (Figure 7.3A)

shows the particle morphology is isotropic and particle size is consistent with

7.1 SiO2 Shells

77

the calculated crystal size. The particles are highly crystalline as observed

from the lattice fringes from HR-TEM. After silica coating the particles

remain isotropic in shape, and little difference is observed from the TEM

images (Figure 7.3B). From HR-TEM a crystalline core with visible lattice

fringes can be seen to be surrounded by a thin shell of a less strongly

absorbing material.

Figure 7.3 | A) TEM of γ-Fe2O3 synthesized in the primary reactor prior to coating.

Particles are isotropic and of narrow size distribution. The insert (HR-TEM) shows clearly

defined lattice fringes. B) TEM of core-shell γ-Fe2O3@SiO2. Insert (HR-TEM) shows the thin

SiO2 shell surrounding the iron oxide core. The white arrow indicates the shell edge and the

black arrow indicate core/shell interface.

STEM-EDX (Figure 7.4A) shows the γ-Fe2O3@SiO2 as a core-shell

nanocomposite. The HAADF detector offers greater contrast in electron

density, but even with this technique the silica shell is not easily detected.

From STEM-EDX (Figure 7.4B-D) the SiO2 particle is 4 nm larger than the γ-

Fe2O3 core particle, and the shell thickness becomes approximately 2 nm or

less. The Si/Fe molar ratio of the product was 0.1 as determined by XRF and

most of the TEOS precursor is therefore unreacted. The product pH was

approximately 6.

Chapter 7 Silica

78

Figure 7.4 | STEM-EDX of γ-Fe2O3@SiO2 from AIC and TEOS at 100°C. γ-Fe2O3@SiO2

adopts a core-shell morphology of 2 nm shell thickness.

7.1.2 TiO2@SiO2

The titania nanoparticles were identified from PXRD (Figure 7.2B) as phase

pure anatase nanoparticles. Structural analysis gave calculated average

crystal sizes of 6.5 nm and 7.3 nm for the coated and uncoated sample,

respectively, in good agreement with the TiO2 studies at these flow rates. The

only change in the diffraction profile was the presence of a large amorphous

bump around 2θ=25° from silica. A Si/Ti concentration ratio of 4 was used due

to the high surface/volume ratio of the small anatase nanoparticles, but from

XRF the Si/Ti ratio of the product is only 1.3, as only a third of the TEOS

precursor is converted into silica. TEM (Figure 7.5) shows no changes to the

overall particle shape when silica is deposited. The overall low particle size

and narrow size distribution of single-phase titania is retained for the core-

shell product. STEM-EDX (Figure 7.6) show the silica shell is several times

thinner than the core nanoparticle diameter. The shell is uniform with a

thickness of less than 2 nm. Formation of titania from TTIP does not change

7.1 SiO2 Shells

79

the pH significantly, and the product suspension pH remains around 4 as

determined by the 1 mM acetic acid solvent.

Figure 7.5 | A) TEM of TiO2 synthesized in the primary reactor. B) TEM of TiO2@SiO2.

The narrow size distribution is preserved and isotropic morphology of the core translates

into an isotropic morphology of the core-shell due to the thin and uniform silica shell.

Figure 7.6 | STEM-EDX of TiO2@SiO2. The elemental maps show thin silica shells of

approximately 1 nm thickness surrounding each titania particle.

Chapter 7 Silica

80

7.1.3 α-Fe2O3@SiO2

Figure 7.2C shows phase pure α-Fe2O3 (ICSD 15840) with calculated average

crystallite sizes of 25.0 nm and 26.0 nm for the SiO2-coated and uncoated

sample, respectively. This confirms that no significant change occurs to the

core nanoparticle in terms of crystal size or structure. A Si/Fe concentration

ratio of 2 was used, however, from XRF the Si/Fe ratio of the product was 3.5

which shows an iron deficit. This is due to hydrolysis of the Fe(NO3)3

precursor which forms nitric acid. This lowers the pH to approximately 2

which partially dissolves the nanoparticles and lowers the yield. TTIP and

AIC are expected to be completely converted into oxides and can therefore be

used as an internal standard for TEOS conversion, but partial dissolution of

the hematite nanoparticles makes this impossible for the α-Fe2O3@SiO2

sample.

Figure 7.7 | A) TEM of α-Fe2O3 prior to coating. Insert (HR-TEM) shows crystalline lattice

fringes. B) TEM of α-Fe2O3@SiO2. The iron oxides are surrounded by silica, but the coating

is incomplete and irregular in shape. Insert (HR-TEM) shows the iron oxides coated in a

lighter material.

7.1 SiO2 Shells

81

Figure 7.8 | A) STEM of α-Fe2O3@SiO2. B-D) STEM-EDX confirms the deposition of silica

on iron oxide.

Figure 7.7 shows the iron oxides are spherical in shape and that the silica

material is highly non uniform and does not form a closed shell. The

thickness of the layer varies which is easily observed at the selected

magnification. Figure 7.8 shows STEM and STEM-EDX of the sample which

confirms the heterogeneous deposition of silica onto iron oxide, and that the

shell coverage is very uneven and varies from incomplete coating and up to 5

nm shell thickness at 100°C. The increase in silica contents from Figure 7.2A-C

can also be observed by the intensity increase of the amorphous bump at 25°.

The presence of SiO2 in all composite samples was confirmed by FTIR and

spectra are listed in Appendix 5.

The influence of temperature and concentration was further investigated for

α-Fe2O3@SiO2. The secondary reactor temperature was increased to 300°C,

but no significant changes in morphology from TEM (Appendix 6) were

observed. This supports that silica growths rapidly at low pH and is likely

fully precipitated long before reaching 300°C. An increased TEOS

Chapter 7 Silica

82

concentration was investigated to see if a more uniform SiO2 shell could be

produced, but this was not the case. STEM-EDX (Appendix 6) reveals the

distribution of multiple iron oxide nanoparticles inside a micron-sized SiO2

agglomerate. This morphology was also observed in 6.3 for the concentrated

Sn precursor in a diluted anatase suspension. The results demonstrate that

shell thickness is not easily tunable with concentration, and high

concentrations of shell relative to core precursor should generally be avoided.

Three different core materials were used, and the three suspensions differ not

only in pH but also in the molecules present (HNO3, NH3, acetic acid, IPA,

citric acid incl. degradation products etc.). A direct comparison is therefore

not straightforward, as many factors differ. Nevertheless, the main reason for

the morphological change in the silica shells is most likely due to pH. The

formation of thin and uniform silica coatings on maghemite and anatase

suggests that formation of individual silica nanoparticles is highly

unfavorable and is only formed due to heterogeneous deposition on the

nanoparticles. After the formation of a few monolayers of silica the

composites have the chemical appearance of silica particles, and further

growth of the shell is arrested. Hematite showed the highest TEOS to SiO2

conversion, as the low pH catalyzes the hydrolysis of the precursor. Only a

third of the TEOS precursor is converted for anatase cores due to the higher

pH, although the high specific surface area of the small particles significantly

contributes to the conversion efficiency. The maghemite suspension is close to

pH neutral, the specific surface area is low, and only 5 % TEOS is converted

into SiO2.

The pH of the anatase nanoparticles was lowered to 4 as this is an

intermediate between the isoelectric points of silica (pH 2) and anatase (pH

6).96 The two materials will therefore be oppositely charged which creates a

strong electrostatic driving force towards heterogeneous deposition. For

maghemite the isoelectric point is 6-7 which is close to the suspension pH, but

this does not seem to affect the uniformity of the shell. This may be related to

the slow silica growth at neutral pH. For hematite the isoelectric point is 8-

7.2 Sub-conclusion

83

9,96 but the low synthesis pH is close to silicas isoelectric point, and the acidic

medium accelerates its nucleation and growth. Silica formation appears to be

initiated at the iron oxide surface, because the nanoparticles all seem to be

partially coated by SiO2. Nevertheless, further growth is fast, and the locally

available TEOS precursor is rapidly consumed by the growing silica instead

of being deposited on the hematite surface. The secondary material is

therefore irregular and non-uniform in shape and does not adopt a core-shell

structure.

7.2 Sub-conclusion

Silica is an important shell material for many applications, and the results

demonstrate that silica coated metal oxides can be readily prepared using the

dual-stage flow reactor. Thin silica shells are important in order to preserve

reactivity in the case of a photocatalytic core, but a thin shell is generally

desirable to preserve the low particle size.

The dual-stage flow reactor is therefore a useful tool, as the limited reaction

time and pH control can restrain silica growth and produce thin silica shells

of 1-2 nm in high yields. Higher reaction temperatures will likely lead to

better conversion of the silica precursor, and this may prove an important

parameter for increasing shell thickness for further studies.

The formation of core-shell morphologies was based on the interplay of pH

and the IEPs of the core and the shell. This concept is further explored in

Chapter 8 to see if suitable synthesis parameters can be selected prior to

synthesis and less by trial-and-error.

85

Chapter 8

IEP and Morphology

The formation of TiO2@SnO2 nanocomposites and core-shell nanoparticles

with silica coatings in previous chapters were linked to the synthesis pH in

the secondary reactor and its value relative the isoelectric points (IEP) of both

core and shell. If the synthesis pH is intermediate of the IEPs of primary and

secondary materials then electrostatic attractions should result in core-shell

formation.

This chapter investigates if IEPs can be used to select a suitable synthesis pH

for composite formation. The predictive power of this IEP concept was tested

with NiO, CuO and ZnO as targeted shell material. The chapter

demonstrates that the core-shell morphology can indeed be induced by choice

of pH, but that the secondary material may differ completely in composition,

crystal size and structure from the reference due to interactions with the

core-nanoparticle suspension. The IEP concept is validated as a useful tool for

selecting suitable synthesis parameters for nanocomposites and is further

applied in Chapter 9 and Chapter 10.

Chapter 8 Core-Shell and IEP

86

8.1 IEP predictions

The different syntheses performed in Chapter 6 and 7 can be divided into

three different composite categories which are illustrated in Figure 8.1. In the

case of a strong electrostatic attraction between core surface and shell

precursor the core is completely coated by a thin shell of uniform thickness.

During shell formation the surface charge changes sign, and this finally stops

the shell growth because the shell precursor is now repelled from the shell.

This was observed for the γ-Fe2O3@SiO2 and TiO2@SiO2 core-shell

nanocomposites in Chapter 6. Contrarily, the two phases will be repelled if

core surface and shell precursor have identical charges as observed for the

alkaline TiO2&SnO2 syntheses. Finally, if either core surface or shell

precursor is neutral then the core may be coated through direct

heterogeneous nucleation as determined by thermodynamics. In that case the

secondary material does not form a closed shell but only decorates the core.

This was observed for the TiO2@SnO2 and α-Fe2O3@SiO2 where the low

synthesis pH was close to the IEP of the shell material. This is likely the best

method for increasing the shell thickness but it is very sensitive to pH and

neutrally charged particles may agglomerate.

The bond distances and charge distributions of the core surface atoms and of

the shell nuclei/complexes are factors which also significantly affect

heterogeneous nucleation. However, these factors may often only be deduced

post-synthesis and holds little predictive power. They are therefore not

further discussed in this dissertation but may be invoked for better

description of core-shell formation from in-situ studies.

The IEP&pH model is a simplistic model which only describes the Coulombic

interactions for composite formation. The IEP is the pH value where the

surface charge is zero, and using it actively provides an important tool for not

only synthesizing nanocomposites but also to select the composite morphology

as illustrated in Figure 8.1A and Figure 8.1C. Three different compounds were

selected as targeted shell materials on a γ-Fe2O3 core: NiO, CuO and ZnO.

8.1 IEP predictions

87

Figure 8.1 | Sketch of formation of composite or mixed particles from electrostatic

interactions. A) The shell precursor is attracted to the oppositely charged core surface. B)

Synthesis pH is below both materials IEP, and a mixed product is formed due to repulsions.

The same result applies if pH is above both IEPs. C) At least one component is

electrostatically neutral. No electrostatic interactions are present. The shell may form only

from direct nucleation, and the secondary material decorates the core non-uniformly.

NiO and CuO are interesting as shell materials, because they are

antiferromagnets. Bi-magnetic core-shell nanostructures comprising a

magnetic core and an antiferromagnetic shell are of great technological

interest, because the two phases are able to interact through the exchange

bias which can enhance the magnetic properties of the composite.97 NiO and

CuO have previously been synthesized hydrothermally by Sue et al.98 in a


88

flow reactor at 400°C. Takami et al.99 synthesized NiO in a flow reactor at

350°C from Ni(NO3) and H2O2.

ZnO is a widely used semiconductor and applications include gas sensors,

varistors, optical devices and catalysis. It often crystallizes as wurtzite and is

readily obtained at even “low” temperatures (~50°C). Extended nanorods are

a common motif but platelets, isotropic particles and other morphologies have

also been reported. Moreover, the role of the alkali cation concentration has

been shown to inhibit the ZnO particle growth by shielding the surface from

the precursor.100-102 Magnetic core and ZnO shell nanocomposites have been

shown to enhance the photocatalytic properties of ZnO compared to phase-

pure ZnO, and in addition a magnetic core provides an easy method for

product collection for recycling of the catalyst.103, 104

The IEPs of the three divalent transition metal oxides are between 9 and 11.

Using maghemite (IEP < 7) as core material a pH range of 7 to 9 was targeted

for synthesis. NH4Cl was selected as buffer because its pKa value is 9.24, and

it is therefore suitable for maintaining a synthesis pH in the required interval.

Maghemite was synthesized from 0.1 M AIC in the primary reactor at 360°C,

and the hot solvent was NH4Cl/NH4OH. Using the buffer as solvent in the

primary reactor stage eliminates the use of an extra pump which increases

residence time. Chlorides were previously discussed as a problem for

hydrothermal synthesis; however, in the present case the buffer solution is

not acidic but weakly alkaline. Initially, NH4NO3 was used as buffer, but the

synthesis was terminated, as it caused the AIC precursor to crystallize into

hematite which is formed from iron nitrate. With the unforeseen emergence

of a red suspension NH4NO3 was therefore replaced with NH4Cl. The

standard (Figure 7.1) dual-stage configuration was employed.

8.2 Copper-based Shell

A copper based shell was prepared by mixing 0.1 M copper nitrate with a

maghemite suspension and heating the mixture in the secondary reactor at

8.2 Copper-based Shell

89

250°C and 350°C (upper and lower heater temperatures). Residence time was

approximately 50-60 seconds, and pH of the collected product was neutral for

all syntheses. TEM (Figure 8.2A) shows the product as isotropic nanoparticles,

and STEM-EDX (Figure 8.2B) reveals that the nanocomposites have clear

core-shell morphology with an iron based core and a copper based shell of

uniform thickness. The shell thickness is approximately an order of

magnitude lower than the core diameter. A thin shell is difficult to detect by

PXRD because of the low volume, and the scattering power of the shell

material is not significantly higher than that of the core. Moreover, the

background is high due to fluorescence from iron, and the X-ray penetration

depth is lowered so the overall diffracting volume is decreased.

The Cu signal from STEM-EDX is very noisy due to sample preparation.

STEM samples are deposited on a copper grid, and this gives a high Cu

background present in all STEM-EDX analyses. From PXRD the diffraction

pattern (Figure 8.3) is dominated by the iron oxide signal, but two small and

wide peaks are seen to emerge at 36.6° and 38.6°.

The peak at 38.6° is most likely the superposition of the (200) and (111)

reflections of monoclinic CuO (ICSD 67850). Copper oxide has its other main

peak for the (002) and (-111) reflections, but that peak coincides with the (311)

main peak of maghemite at 35.6°.

The other emerged wide peak at 36.6° coincides with the (221) peak of

maghemite, but is nevertheless most likely related to the copper compound

because it is much wider than the maghemite peaks. This peak is probably

the (111) main peak of Cu2O. This compound is formed from reduction of

Cu(II) to Cu(I). Deng et al.105 showed that Cu2O impurities formed during

reduction of Cu(II) to Cu(0) in the presence of citric acid, and that this

impurity was absent when using other carboxylic capping agents in an

alkaline NH4OH-based medium. The authors furthermore suggested that a

reduction in particle size was due to formation of copper(II)-ammonia

complexes which stabilized the particle surface and limited the growth rate.


90

Accurate size determination of the surface bound CuO and Cu2O is not

possible due to the low crystalline content and iron fluorescence. Furthermore,

peak width if most likely also affected by strain from lattice mismatch along

the core/shell interface. Additionally, diffusion and substitution reactions

might occur which would cause a peak shift and result in peak widening.

The γ-Fe2O3-free reference yielded CuO nanoparticles of 25 nm, but other

peaks are observed at lower diffraction angles. Cu forms a series of layered

hydroxide structures. Monoclinic Cu2(OH)3NO3 (ICSD 31353, a = 5.61 Å, b =

6.09 Å, c = 6.93 Å, β = 94.48º) was added to fit the data, and Figure 8.3 shows

the diffraction data modelled with both CuO and Cu2(OH)3NO3. The large

peak at 11.7º is then the (002) reflection of Cu2(OH)3NO3, but the c-axis is

expanded 9% to 7.55 Å to fit the data. This fit also provides a good description

of the peak at 24º, but the large c-axis expansion suggests a layer-by-layer

distance greater than that of pure Cu2(OH)3NO3. This may be due to the

incorporation of different anions as previously demonstrated.106 The Cl/Cu

molar content was 1%, and the intercalated specie is therefore not chlorine. A

single, sharp peak at 16.2º could not be identified, but could be from a CuO

supercell structure.

Figure 8.2 | A) TEM of γ-Fe2O3@CuxO. Particles are isotropic and surface is smooth. B)

STEM-EDX of isotropic γ-Fe2O3@CuxO nanoparticles. High Cu-background due to STEM

copper grids.

8.3 Nickel-based Shell

91

Figure 8.3 | PXRD of Cu(II) reference (blue, top) and composite (red, bottom) products.

The reference yields mainly CuO, but both CuO and Cu2O are present in the composite

product.

Better data quality can be obtained by switching the Rigaku Cu-target to

another anode which would eliminate iron fluorescence and decrease the

background, but synchrotron radiation is the best option for accurate

determination of the structure of the thin shells.


Replacing Cu(NO3)2 with Ni(NO3)2 produced an identical core-shell

morphology as seen in Figure 8.4. Figure 8.5 (red data) shows the diffraction

pattern of the nanocomposite, and the only crystalline phase detected is that

of maghemite. The results are comparable to the γ-Fe2O3@NiO core-shell

hybrids synthesized by Skoropata et al.,107 who demonstrated an increase in

magnetic coercivity compared to phase pure γ-Fe2O3.


92

The maghemite-free nickel reference produced β-Ni(OH)2 (ICSD 169978) and

cubic NiO (ICSD 9866) but also at least one other phase which could not be

identified (Figure 8.5, blue data). The presence of β-Ni(OH)2 is in agreement

with the results by Liang et al.,108 who obtained β-Ni(OH)2 from hydrothermal

treatment at 200°C in an autoclave for 2 hours. They speculated that the

addition of ammonia lead to the formation of Ni(NH3)62+ complexes which

affected the growth rate of nickel hydroxide.

NiO is obtained from dehydration of nickel hydroxide. From PXRD the

intensity of the (111) peak at 37.2° is much higher than the (200) peak at

43.3°. For an isotropic crystal shape the (200) peak is the main peak, and the

high intensity of the (111) peak indicates that the crystal shape is highly

anisotropic, which is not expected for a cubic crystal structure. This is likely

due to the heating profile in the secondary reactor. Particles of β-Ni(OH)2 are

quickly precipitated, but dehydration into NiO is either slow or does not

proceed until further down the reactor, where the temperature is high enough.

The hexagonal β-Ni(OH)2 particle shape is most likely platelets, and this

morphology is preserved when the nickel dihydroxide is converted into the

oxide. This observation was also made by Takami et al.,99 who synthesized β-

Ni(OH)2 platelets in a hydrothermal flow reactor and obtained NiO by

calcination of the product. The same relative peak intensities were also

obtained in that study and shown by SEM to be NiO platelets.

While only the maghemite phase is present in the diffraction pattern in

Figure 8.5 (red data) its diffraction profile has changed considerably. For

instance the intensity of the (004) reflection at 43° has increased and become

the second largest peak in intensity – it is usually only the fourth largest.

Moreover, the average γ-Fe2O3 crystal size has increased to 20.0(1) nm. The

maghemite core has been synthesized using identical conditions for every

synthesis and is usually around 17-18 nm. The Ni-based shells were

synthesized during the same synthesis run as the Cu-based shells, and the

maghemite crystal size of the γ-Fe2O3@CuO/Cu2O nanocomposites was 18.6

nm. These discrepancies may be related to the nickel coating.


93

Figure 8.4 | A) STEM of maghemite nanoparticles with a Ni(II)-based shell. B-D) EDX

shows a uniform shell of a few nanometers surrounding the iron oxide core.

Figure 8.5 | PXRD of nickel reference (blue, top) and composite (red, bottom) products.

Only the core iron oxide diffraction signal is observed for the hybrid. The shell composition

could not be accurately identified, but is a nickel(II) compound and is denoted NiOxLy.


94

The crystal size is usually smaller than the particle size, because the particle

consists of a crystalline core and an amorphous shell. It is possible that a

nickel compound is grown epitaxially on the maghemite surface, which could

move the amorphous shell into the nickel layer and fully crystallize the iron

oxide particle. Nickel is closer to iron than copper in the periodic table and

should therefore a priori be better at epitaxial growth than copper which

would induce considerable strain in the outer core. The cubic unit cell of

spinel structured γ-Fe2O3 is 8.39 Å while that of NiO is 4.177 Å. The lattice

mismatch of a γ-Fe2O3 unit cell length compared to two NiO unit cells is

therefore only 0.5%, and it is possible that NiO grows epitaxially from the

maghemite surface. From TEM analysis (Figure 8.6) the surface of the

nanocomposites are rugged and not as smooth as the CuO/Cu2O shells in

Figure 8.2 or phase-pure maghemite (Figure 7.3A). From the above

observations it may be speculated that the shell material is NiO, but the

dataset cannot prove this.

Figure 8.6 | TEM of γ-Fe2O3@NiOxLy core-shell nanocomposites. The overall particle

shape is isotropic, but the surface is more rugged and uneven compared to phase pure

maghemite.

8.4 ZnO and Zn6(OH)6(CO3)2

95


0.1 M Zn(Ac)2 was used as secondary precursor. Both precursors were

pumped at 5 mL/min while the hot solvent flow rate was 15 mL/min. The

solvent was a 0.5 M NH4OH and 0.3 M NH4Cl buffer. Neutralization of the

zinc precursor is usually necessary, as the acid formed by hydrolysis will

redissolve any particles that may have precipitated, and the alkaline buffer

serves a dual purpose in this case.

Figure 8.7 shows the diffraction patterns of the reference and the composite

product using the 0.5/0.3 M buffer at 250°C. Wurtzite structured zinc oxide

(ICSD 34477) is formed in the absence of maghemite and the average crystal

size 36 nm. Contrarily, when maghemite is present the large ZnO particles

are no longer formed and the precursor crystallizes into smaller

Zn6(OH)6(CO3)2 (hydrozincite, ICSD 16583) of 7 nm crystallite size. This

phase is formed in response to high levels of CO2 from the degradation of

ammonium iron citrate and is an intermediate phase when converting ZnO to

ZnCO3.109 Hydrozincite is one of several layered zinc-hydroxide structures in

which several anions may be introduced including acetates and nitrates110, 111.

Chlorine is present in the medium as a NH4Cl buffer, but no chlorides were

present in the solid product as determined by XRF.

To investigate if the change in crystal growth is only due to the presence of

carbonates or also due to the iron oxide nanoparticles the product was

investigated with STEM. Figure 8.8 shows the product as isotropic

nanoparticles, and the particle size correlates with that of the maghemite

crystal size. No smaller individual particles are seen as would otherwise be

expected if hydrozincite nanoparticles were ungrafted to the iron oxides.


96

Figure 8.7 | PXRD of ZnO reference (blue, top) and γ-Fe2O3@Zn6(OH)6(CO3)2 composite

(red, bottom) products. The maghemite suspension contains carbonates which promotes

formation of Zn6(OH)6(CO3)2 over ZnO.

Figure 8.8 | STEM of γ-Fe2O3@Zn6(OH)6(CO3)2. The nanoparticles are isotropic and the

particle size is in agreement with the maghemite crystal size.


97

STEM-EDX (Figure 8.9) shows a thin and uniform zinc-based shell around the

iron oxide particles. The morphology is a core-shell structure with uniform

shell thickness of approximately 2 nm. Shell thickness was estimated by

comparison of core and shell diameters from EDX-maps. By comparing the

Fe/Zn molar ratio obtained from XRF with the calculated crystalline weight

fractions from Rietveld analysis only 31 wt% of the zinc in the product is

crystalline assuming full maghemite crystallinity. This explains the

discrepancy between the thin shell and larger hydrozincite crystal size. A

single zinc-based particle can be seen on the shell which has a particle size

identical to the calculated hydrozincite crystal size.

Figure 8.9 | STEM-EDX of γ-Fe2O3@Zn6(OH)6(CO3)2. The iron oxide is coated in a thin

zinc-based shell. A clear void is observed in (D) from where the core sits.


98

Calcination of the product should convert the hydrozincite into ZnO, but this

would also lead to diffusion across the interface and formation of mixed iron-

zinc oxides. The temperature dependence of this composite formation was

further investigated with a weaker solvent buffer of 0.08 M NH4OH and 0.08

M NH4Cl. The temperature was varied from 85°C to 350°C, but the crystal

size of hydrozincite remained fixed at 4.0(5) nm for all temperatures. This is a

decrease in crystal size compared to using the stronger buffer. No significant

change in crystalline weight fractions from Rietveld analysis was observed.

The constant crystal size with reactor temperature suggests that the

hydrozincite is rapidly grown on the core particles. If the zinc species are

immobilized on the maghemite surface no precursor is available for further

crystal growth even at elevated temperatures. Hydrozincite crystallization is

observed above 85°C, but nanocomposite formation could potentially be

completed before even entering the secondary reactor.

It is interesting to note that all maghemite-free (and carbonate-free)

references yielded phase pure ZnO. The average ZnO crystal size increased

with temperature from 17 nm (85°C) to 34 nm (250°C) and finally 50 nm at

350°C. At an intermediate synthesis of 85°C in the top reactor heater and

250°C in the bottom reactor heater the average ZnO crystal size was 31 nm.

These values correspond well with those obtained by Søndergaard et al.100

The structural model used in the Rietveld refinements assumes an isotropic

particle shape which is not the case for the ZnO nanorods. The data

demonstrates a trend of increasing crystal size with temperature but does not

mandate a quantitative interpretation, much like the extended SnO2

nanorods at low pH discussed in Chapter 6.

Lowering of the maghemite precursor concentration from 0.1 M to 0.02 M

decreased the carbonate concentration and increases the Zn/Fe molar ratio

from 1 to 5. This increased the hydrozincite crystal size by up to 50%, but also

led to the formation of large ungrafted ZnO nanorods coexisting with the

8.5 Sub-conclusion

99

hydrozincite shells (Figure 8.10). This demonstrates that ZnO cannot be

deposited on hydrozincite at the applied reaction pH.

Figure 8.10 | STEM-EDX of isotropic γ-Fe2O3 nanoparticles and ZnO nanorods. Increasing

the Zn concentration lead to ungrafted ZnO nanorods.

The isoelectric point of hydrozincite is below 6, and the nanocomposite

surfaces are negatively charged.112 From IEP considerations composite

formation should not be expected to occur if the shell is amorphous

hydrozincite. In the study of Bitenc et al.113 hydrozincite is shown to form

from condensation reactions of the neutral complexes [Zn(OH)(HCO3)(OH2)3]

and [Zn(OH)2(OH2)4] (HCO3 is a bidentate ligand). Composite formation is

therefore more likely due to thermodynamics or the charge distribution of

these complexes and the iron oxide surface.

A more alkaline buffer of 0.3 M NH4OH and 0.1 M NH4Cl (product pH 9.5)

lead to the formation of large ZnO particles with product characteristics

independent of the presence of maghemite and carbonates indicating that

ZnO is ungrafted. No formation of hydrozincite was observed at this elevated

pH. With these observations it was evident that γ-Fe2O3@ZnO cannot be

synthesized from the organic AIP precursor.

8.5 Sub-conclusion

For all three composites a clear core-shell structure was obtained and the

secondary precursor was shown to behave completely different in the


100

presence of the maghemite suspension in all three cases. Little long-range

order exists in the thin shells and the composition of the shell is difficult to

determine with diffraction methods used. Further heat treatment would allow

the shells to crystallize, but this would also promote the formation of an

Fe(III)/Me(II) intermediate layer between the core and the shell. Additionally,

this new intermediate layer might consume the shell layer as iron diffuses

from the core to the surface. Sintering could therefore be used for the

formation of a range of interesting γ-Fe2O3@MeFe2O4 compounds for magnetic

applications. Further treatment of the product is, nonetheless, beyond the

exploration of the IEP concept.

The IEP concept was used in this chapter to synthesize isotropic core-shell.

The shells were uniform and approximately 2 nm in thickness as was also

observed for the silica shells deposited on titania and maghemite in Chapter 7.

The driving force of formation for such composites is likely due to electrostatic

interactions which attracts the secondary precursor to the oppositely charged

core until the particle is fully coated by at least one monolayer of shell

material. A core-shell structure was not observed for the TiO2@SnO2 hybrids

but that is not expected on the basis of IEP alone. Here the driving force

seems to be of a more thermodynamic nature which favors heterogeneous

deposition rather than overcoming the large energy barrier associated with

homogeneous nucleation.

The chapter further demonstrates that magnetic core-shell nanoparticles can

be readily synthesized hydrothermally in the dual-stage reactor. Further

studies may look into control of the shell material and especially on how to

increase the thickness of the shell without producing ungrafted particles or

large agglomerates. Synthesis pH is likely to influence this, but the identity

and concentration of the precursor are also factors to be considered. This is

the topic for the following chapter.

101

Chapter 9

γ-Fe2O3@TiO2

Isotropic core-shell structures were synthesized in Chapter 7 and Chapter 8

with shells of only a few nanometers. This chapter investigates how shell

thickness can be increased with γ-Fe2O3@TiO2 as an example. In Chapter 6

SnCl4 and Sn(NO3)4 demonstrated that different precursors act differently

towards composite formation, and the role of the shell precursor is further

investigated in this chapter.

TiO2 is a highly interesting material by its own right, but it can also be

considered an intermediate towards BaTiO3 and SrTiO3. This chapter

describes how different precursors and concentrations influence synthesis of

γ-Fe2O3@TiO2, but the ultimate goal is to lay the foundations for multiferroic

γ-Fe2O3@BaTiO3 and γ-Fe2O3@SrTiO3 in Chapter 10.

Chapter 9 γ-Fe2O3@TiO2

102

9.1 Magnetic photocatalysts

With the extensive synthesis work already performed using both metal oxides

as core-material it is a natural step to investigate their hybrid. The initial

motivation for a γ-Fe2O3@TiO2 core-shell synthesis was therefore not for

applications but to gain a deeper insight into the formation chemistry of

heterostructures using these well-known compounds. γ-Fe2O3@TiO2 hybrids

are actually highly unsuitable for applications. While deposition of a

catalytically active material on a magnetic core allows for easy collection e.g.

for recycling this necessitates that the performance is not lowered compared

to the unsupported catalyst. This is unfortunately the case for γ-Fe2O3@TiO2.

The photocatalytic activity of titania is reduced when deposited on a magnetic

core because the electronic interactions across the interface leads to an

increase in electron-hole recombination. This can be solved by insertion of a

sandwiched SiO2 layer which also prevents photodissolution of iron. TiO2 has

been used as coatings for several different magnetic cores, but often multiple

core particles are agglomerated and covered by larger lumps of titania and

the composites are often in excess of 100 nm.114-116 The hybrid has also been

employed for photokilling of bacteria.117

9.2 The influence of precursor and concentration

The deposition of a titania layer on maghemite was performed using the

standard dual-stage configuration (Figure 7.1). The synthesis was performed

initially with TTIP as Ti precursor. This was in retrospect a futile experiment.

The TTIP precursor reacted immediately with the water in the aqueous

maghemite suspension during cold mixing. The collected product was a

darker grey and not as black as pure phase pure maghemite. Maghemite

typically sediments within a few seconds of centrifugation while titania

requires several minutes at 6500 rpm. After a few seconds the dark grey

suspension had changed to almost complete white while a dark precipitate

was observed at the bottom of the tube. A two phase mixture or composite


103

particles could therefore be visually identified following a few seconds of

centrifugation. XRF confirmed only trace amounts of titanium in the black

powder. After extended centrifugation all of the TiO2 sedimented which

formed a white layer on top of the black iron oxide.

The non-aqueous Ti-precursor, Titanium tetraisopropoxide (TTIP), was

therefore replaced with the aqueous Ti-precursor Titanium bis(Ammonium

Lactato) Dihydroxide (TALD), for TiO2 formation. No attempts were done to

regulate the pH as the maghemite suspension is approximately 6 following

hydrolysis of ammonium iron citrate which is close to the IEPs of both oxides

(IEP of anatase 4-8, IEP of γ-Fe2O3 3.5-6.5, Figure 6.9). Centrifugation of the

composites yielded a dark grey precipitate after only a few seconds while the

near-white Ø&TiO2 references required up to 10 minutes to completely

sediment. These two different two-phase products are shown in Figure 9.1 (left

and right) along with the suspended product of one of the samples in the

middle.

Figure 9.1 | Determination of a composite or mixed powder can sometimes be seen

directly from centrifugation. A dark grey suspension (middle) may sediment into a dark

grey precipitate (left) or a black and white sediment (right). TALD produces a composite of

iron oxide and titania (left) whereas TTIP produces a mixed product (right). TiO2 forms a

layer on top of the iron oxide because it takes longer to sediment.


104

Figure 9.2 shows the diffraction patterns for γ-Fe2O3&TiO2 and the

corresponding Ø&TiO2 references synthesized at 250°C from TALD. PXRD of

the products obtained at 350°C are included in Appendix 7. At 350°C the

secondary reactor temperature is equal to the primary reactor temperature,

and this causes the maghemite crystal size to increase from 17.6 nm (250°C)

to 18.1 nm (350°C).

Figure 9.2 | PXRD of γ-Fe2O3@TiO2 at 250°C with a Ti:Fe ratio of 1 (blue, middle) and 2

(green, bottom). Reference synthesis yielded phase pure anatase (red, top).

When maghemite is present in the stream the average anatase crystal size

decreases from 3.2 nm to 2.6 nm at 250°C and from 5.9 nm to 4.7 nm at 350°C.

Nevertheless, a 25% reduction in anatase crystal size is not enough to

conclude that maghemite induces a significant change in TiO2 formation and

certainly not based on only two different temperatures.

STEM analysis shows that TiO2 is indeed deposited heterogeneously on the

iron oxides. Figure 9.3 shows multiple smaller anatase particles surrounding


105

the iron oxides cores at two different temperatures. The best coverage is

observed at 250°C, while at 350°C a large fraction of the titania nanoparticles

are ungrafted. In addition, the larger anatase crystal size at higher

temperatures will further decrease the iron oxide coverage, and the higher

temperature is therefore to be avoided.

Figure 9.3 | TEM of γ-Fe2O3@TiO2 with a Ti:Fe molar ratio of 1. A better coverage is

obtained at 250°C (A) rather than at 350°C (B). More titania is ungrafted at higher reaction

temperatures.

STEM-EDX (Figure 9.4) confirms that the nanocomposites are core-shell

structures, although the shell is not as uniform as has been observed in

previous chapters. This is likely due to the synthesis pH which is close to 6,

and the deposition mechanism is not due to electrostatics. The maximum

shell thickness is consistent with the crystal size of the secondary phase much

like the TiO2@SnO2 nanocomposites in Chapter 6. But unlike the TiO2@SnO2

hybrids a better coating of the core is not observed with higher temperatures

for γ-Fe2O3@TiO2. This may be related to differences in reaction rates and

mechanisms for the two systems. It is likely that an even more uniform

anatase shell can be obtained at lower temperatures. Further studies may

elucidate the changes to the shell morphology at different temperatures.


106

Figure 9.4 | STEM-EDX of γ-Fe2O3@TiO2 at 250°C with a Ti:Fe molar ratio of 1.

From these observations a more complete coverage and closed shell

morphology would therefore be expected for a higher Ti/Fe ratio. This was

verified by doubling the TALD concentration. TEM (Figure 9.5) confirms that

the product is still a hybrid and that only a very small fraction of the titania

nanoparticles are ungrafted. The shell layer is denser but not significantly

thicker (Figure 9.6). With a doubled TALD concentration the calculated

average crystal size was 2.5 nm which indicates that TALD concentration

does not affect the crystal size as was also observed for the TTIP precursor in

Chapter 5. From XRF the Ti/Fe molar ratio of the product was 1.9.

The reverse hybrid was also attempted for a TiO2@γ-Fe2O3 nanocomposite,

but the product (not shown) was made up of mixed particles. As previously

mentioned in 6.6 no composites were ever obtained when an anatase core was

synthesized from TALD. TTIP may yield different results although the

structure is mainly of interest in order to understand the deposition

mechanism. Deposition of a black material on a photocatalyst is of limited use

in applications.

9.3 Sub-conclusion

107

Figure 9.5 | TEM of γ-Fe2O3@TiO2 at 250°C with a Ti:Fe molar ratio of 2. The increase in

Ti concentration does not lead to an increase in ungrafted titania particles.

Figure 9.6 | STEM-EDX of γ-Fe2O3@TiO2 at 250°C with a Ti:Fe molar ratio of 2.

9.3 Sub-conclusion

Core-shell γ-Fe2O3@TiO2 nanoparticles were successfully synthesized at near

neutral pH. The anatase shell becomes denser and probably also thicker with

increasing Ti content, and a reaction temperature of 250°C or less is optimal.


108

The shell is less uniform than those reported in Chapter 7 and Chapter 8. The

IEPs of both core and shell were close to the synthesis pH and electrostatic

attractions are not the main driving force for composite formation. The titania

shell growth was therefore not arrested from repulsion with the precursor

species as observed for γ-Fe2O3@SiO2. Instead the titania shell continued to

grow until the local precursor concentration was depleted, as observed from

an increase in Ti content with higher Ti concentration.

The results provide important information on how to synthesize γ-

Fe2O3@TiO2 as a core-shell nanocomposite. This is the intermediate in the

synthesis of γ-Fe2O3@SrTiO3 and γ-Fe2O3@BaTiO3 discussed in Chapter 10.

109

Chapter 10

Multiferroics

In the previous chapters several different shell materials have been

investigated: From amorphous silica to crystalline metal oxides in different

oxidation states. Although the chemistry is complex the synthesis pathway

has been practically the same: Mixing a secondary precursor with a

suspension and leading the mixture into a secondary reactor for formation of

the secondary phase.

This chapter deviates from this synthesis route and demonstrates how

combination of reactor configuration and chemical know-how can produce

some of the most advanced materials of present day.

Chapter 10 Multiferroics

110

10.1 BaTiO3 and SrTiO3 Ferroelectrics

The dual-stage reactor has been shown as a highly versatile apparatus for the

synthesis of nanocomposites. However, composite formation is dependent on

pH and many compounds are not formed at the pH required for

heterogeneous deposition. BaTiO3 and SrTiO3 are examples of such

compounds which only precipitate at high pH. Mixing γ-Fe2O3 nanoparticles

with Titanium bis(Ammonium Lactato) Dihydroxide (TALD), Ba(NO3)2 and

NaOH may lead to γ-Fe2O3 and BaTiO3 but never as a single-particle

composite material under the highly alkaline conditions. The synthesis

strategy was therefore not to deposit BaTiO3 onto maghemite but first form a

γ-Fe2O3@TiO2 composite and subsequently convert TiO2 into BaTiO3.

Hydrothermal conversion of TiO2 into BaTiO3 has previously been

demonstrated.118, 119 The reaction strategy is illustrated in Figure 10.1.

Figure 10.1 | Illustration of initial synthesis strategy to obtain multiferroic γ-

Fe2O3@BaTiO3. Magnetic nanoparticles synthesized in the primary reactor stage (not

shown) are mixed with Ti4+ and Sr2+ reactants for formation of TiO2 (Step 1). Unreacted

Sr2+ converts TiO2 into SrTiO3 when the pH is raised by NaOH addition in the final heater

(Step 2). NaOH is preheated to maintain a stable reaction temperature.

BaTiO3 and SrTiO3 are both ferroelectrics. They crystallize in the perovskite

structure, and in the cubic phase the Ti atoms sit in the center of a regular

octahedron. With the application of an external electrical field the Ti atoms

10.1 BaTiO3 and SrTiO3 Ferroelectrics

111

can be displaced from their octahedral equilibrium position, and the crystal

symmetry is reduced from cubic to tetragonal.

When the external field is removed the TiO6 octahedra remain distorted, and

the compound possesses a net polarization. Switching the orientation of the

field will then induce a hysteresis loop. This tetragonal distortion can be

quantified by the c/a ratio, and an increase in c/a ratio correlates with higher

ferroelectric polarization. BaTiO3 is tetragonal (ICSD 67519, a = b = 3.9998 Å,

c = 4.0180 Å, c/a = 1.0046) at room temperature and has a spontaneous

polarization, but becomes cubic (ICSD 67518, a = b = c = 4.006 Å) above the

Curie temperature, Tc, at 120°C. This temperature is, however, lowered with

crystal size, and for crystals smaller than 100 nm the cubic phase is the most

stable at room temperature. The Curie temperature is further lowered by

substituting Ba with Sr or with the application of an external pressure.120, 121

When ferroelectricity is combined with ferromagnetism the compound is

known as a multiferroic. However, the two phenomena are often mutually

exclusive and usually cannot be present in the same phase, because the

displaced cation in ferroelectrics requires empty d-orbitals whereas the

magnetic moment is usually due to d-electrons. BiFeO3 is a rare example of a

single-phase multiferroic where ferroelectricity is due to the lone pair on the

bismuth atom. Heterostructures of ferroelectric and magnetic phases have

been proposed as the most suitable pathway, and thin-films have been shown

as a versatile method for bringing the two phases into contact with each

other.122 It should be noted that maghemite is a ferrimagnet.

Core-shell particles are interesting in this respect as the morphology

maximizes the interface between the two phases which optimizes the

magnetoelectric interactions. Multiferroic core-shell particles previously

reported are often in excess of 0.1 μm in diameter. Koo et al.123 synthesized

BaTiO3@Fe3O4 core-shell particles from commercial BaTiO3 and FeCl2 via a

sonochemical route. Composites of BaTiO3 and NiFe2O4 composites were

prepared by Sreenivasulu et al.124 using organic linkers on premade particles.

Fe3O4@PbTiO3 and other core-shell multiferroics were prepared


112

hydrothermally by Liu et al.125 who showed considerable strain in the

tetragonal phase from lattice mismatch with the magnetic core. This was also

observed by Corral-Flores et al.126 who showed an increase in tetragonal

BaTiO3 content when deposited on CoFe2O4 core-nanoparticles due to strain,

and that this was related to the thickness of the shell. The core nanoparticle

can thereby be used to stabilize the tetragonal phase which turns the shell

ferroelectric, and the combined nanocomposite becomes a multiferroic. This

has been supported by theoretical calculations which show that

ferroelectricity may indeed be present at room-temperature for perovskite

thin films deposited on a substrate for a thin film thickness as low as 3 nm.127

Multiferroic heterostructures thereby provides a highly interesting case of the

interplay between structure and property: a magnetic core may be used to

induce structural changes in the shell which produces a ferroelectric response

which in return interacts with the magnetic moment of the core.

Formation of a multiferroic compound is a highly complex process with

several parameters including reactor configuration, temperature (gradient)

not to mention the chemicals involved. A methodical course involving the

change of a single parameter at a time would accurately map out the

parameter space, but the number of samples and time required simply makes

this impossible. The results in the following text are presented in a sequence

so as to construct a narrative that reflects the synthesis approach.

10.2 Reactor Reconfiguration

For nanocomposite formation the 110 cm long secondary reactor was replaced

by two 50 cm reactor tubes connected by a mixing tee into which hot NaOH

was mixed. Incidentally, this configuration is identical with the “direct

approach” which was discussed in Chapter 3. The mixing is far from optimal,

but γ-Fe2O3@TiO2 composite formation should be fully completed by this stage,

and the role of the base is only to provide a sufficiently alkaline environment

for incorporation of barium or strontium into the already grafted titania

10.3 BaTiO3 Precursors

113

nanoparticles. The local pH may vary, but this should be insignificant as long

as NaOH is added in surplus. Hydroxides of barium and strontium

precipitate at high pH, and it is therefore not possible to dissolve these salts

directly in the NaOH solvent.

Alternatively, a Ba/Sr salt could be injected by a separate pump. The last

mixing tee between step 1 and step 2 would then need replacing with a

mixing cross. However, this mixing is terribly inefficient due to gradients in

density, composition and temperature much like mixing in the “direct

approach” if not worse.

The suggested synthesis route in Figure 10.1 involves several steps and initial

experiments were performed to test the feasibility of the synthesis plan. A

BaTiO3 reference was initially synthesized without iron oxide present to

verify if the reaction conditions, and especially if the limited residence time

are sufficient to convert TiO2 into BaTiO3.


A recurring theme throughout this dissertation is the importance of precursor

choice for heterogeneous deposition of a material. While the addition of

barium or strontium salts should have little significance for titania formation

they may affect the shell formation. Different precursors were therefore

screened for conversion of anatase into MTiO3 perovskites. All precursor flow

rates were 5 mL/min, 15 mL/min for the primary solvent (H2O) and 10

mL/min for the NaOH solutions.

10.3.1 Ba(OH)2

BaTiO3 without iron oxides

Ba(OH)2 was initially investigated as barium source. HNO3 was added to

neutralize pH of the mixed 0.15 M Ba(OH)2 and 0.1 M TALD precursor. The


114

diffraction pattern (Figure 10.2A.I) shows the formation of 24.6 nm cubic

BaTiO3 at 150°C and 325°C (upper and lower secondary reactor temperatures,

respectively). This was in the absence of maghemite and with a NaOH

concentration of 0.5 M. A minor BaCO3 impurity is detected from residual

carbonates present from degradation of citrates still present in the system.

The precipitation of BaCO3 is another reason for the addition of a Ba surplus

relative to Ti. No peak splitting was observed, and the BaTiO3 phase is cubic

and not tetragonal.

Substitution of TALD with TTIP produced BaTiO3 nanocrystals of 31.1(1) nm

using the cubic model (a = 4.02710(4) Å, RBragg= 11.1 and RF = 5.89). The

diffraction pattern is shown in (Figure 10.2A.III). The BaTiO3 structure is

known from literature to be cubic below 100 nm at room temperature, but a

better fit is expected when switching to the tetragonal structure as more

parameters are available to fit the data. Using the tetragonal model increases

the crystal size to 35.3(1) nm, because the broad peaks are now described as a

peak split i.e. superposition of two narrower peaks. This model improves the

R factors (RBragg = 4.87, RF = 3.44), and the c/a ratio is 4.0334 Å/4.0221 Å =

1.0028.

Despite the better tetragonal fit, a cubic model was selected for all

investigations. The PXRD data quality does not justify using the tetragonal

structure as model because multiple phases are present and iron fluoresces

strongly. Synchrotron radiation or measurement of the dielectric properties

would provide an accurate determination of the phase, but this is beyond the

scope of this dissertation.

The crystal sizes are in agreement with Reverón et al.,128 who obtained cubic

BaTiO3 from the metal isopropoxides in a supercritical flow reactor. They

showed an increase in crystal size from 20 nm to 35 nm when the Ba:Ti molar

ratio was increased from 1 to 2. A small BaCO3 impurity was also detected,

and this phase will be addressed later in the text. It is especially interesting

to note that this rapid heating compared to the more gradual heating in the


115

dual-stage secondary reactor does not seem to change the average BaTiO3

crystal size.

TTIP is unsuitable for heterogeneous anatase deposition as already discussed

in Chapter 9 and is only described here for initial comparison with the TALD

precursor. TTIP hydrolyses in the presence of water, and the synthesis was

performed by pumping 0.1 M TTIP in IPA from a separate pump and mixing

it with aqueous Ba(OH)2 and HNO3 from another pump (five pumps in total).

This mixture was then mixed with the main stream from the primary reactor.

These tests verified that crystalline BaTiO3 can indeed be formed in the

secondary reactor, and the perovskite is obtained at 325°C or below. The

synthesis was therefore continued with the addition of a maghemite

suspension.

BaTiO3 with iron oxides

PXRD of the product obtained when mixing the maghemite suspension from

AIC with a secondary precursor of TALD, Ba(OH)2 and HNO3.is shown in

Figure 10.2A.II. The diffraction pattern shows strong peaks of BaCO3 from

citrate degradation of the Ammonium Iron Citrate precursor. More

importantly the expected maghemite peaks have been replaced by hematite.

This was also observed using TTIP as anatase precursor (Figure 10.2A.IV).

Ba(OH)2 is a strong base and is difficult to neutralize without careful titration.

It is likely that the pH was too low and too much HNO3 was added causing

the maghemite to partially dissolve and recrystallize into α-Fe2O3. The

hematite peak width is identical to the instrumental contribution to peak

width, and the crystals are in excess of 100 nm. This was confirmed by STEM

(Figure 10.2B). Hematite is likely formed from partial dissolution of the

maghemite nanoparticles which then agglomerate and crystallize into

micron-sized hematite particles.


116

Figure 10.2 | A) PXRD of BaTiO3 synthesized in the secondary reactor at 150°C (upper

heater) and 325°C (lower heater) from different Ti precursors. BaCO3 was identified for all

samples. Y-values plotted as square-root of intensities I,II) BaTiO3 synthesized from TALD.

III,IV) BaTiO3 synthesized from TTIP. II,IV) BaTiO3 synthesized in an iron oxide

suspension. B) Iron oxide is converted from maghemite nanocrystals into large hematite

particles due to excess HNO3.

10.3.2 Ba(NO3)2

Ba(NO3)2 is the salt of a strong base and strong acid. The salt was dissolved

in 0.01 mM HNO3 for a final pH of approximately 5 as the dissolution of

Ba(NO3)2 should not change the precursor pH. The PXRD (Figure 10.3A.I)

show the presence of γ-Fe2O3, BaTiO3 (average crystal size of 50 nm) and

BaCO3 as expected. A SrTiO3 crystal size of 27.3 nm were obtained when

using Sr(NO3)2 (Figure 10.3A.II). Ba and Sr forms solid solutions in the BaxSr1-

xTiO3 system, and smaller crystal sizes are obtained for Sr in agreement with

previous reports.129 Nevertheless, STEM (Figure 10.3B) shows that the product

is not a nanohybrid.

It is surprising that the titania is not grafted on the iron oxides. The presence

of nitrates and/or strontium seems to play a role in the titania deposition on

the iron oxides. The strong nitric acid is a poor buffer, and pH may be

unstable. Or maybe titania and metal nitrates are simply a poor combination

for unknown reasons in terms of composite formation as previously observed

for Sn(NO3)4. Whatever the reason, better results were obtained with acetates

10.4 The Two Synthesis Steps

117

and further experiments were performed using this precursor. The pKa value

of acetic acid is 4.76, which should make it more suitable for stabilizing the

pH near this value. Acetic acid has previously been demonstrated as a

chelating agent for Sr and Ti precursors used for SrTiO3 precipitation.130

Figure 10.3 | A) PXRD of MTiO3, MCO3 and γ-Fe2O3 using TALD and M(NO3)2. I) BaTiO3

from Ba(NO3)2 II) SrTiO3 from Sr(NO3)2. B) STEM-EDX of SrCO3, SrTiO3 and γ-Fe2O3.

Large SrCO3 particles are dispersed between ungrafted titania and iron oxide.


In the previous sections it was shown that BaTiO3 can indeed form in the

secondary reactor, but that the use of nitrates interferes with titania

deposition. Synthesis of γ-Fe2O3@TiO2 was discussed in Chapter 9, but for γ-

Fe2O3@BaTiO3 only the first half of the secondary reactor is available for

titania formation, because the lower half is for TiO2 to BaTiO3 conversion.

This section investigates the two steps in Figure 10.1 in order to understand

the chemistry and to make the secondary reactor less of a black box.


118

10.4.1 Step 1: Titania deposition

The TALD precursor was selected for TiO2 formation as TTIP leads to

ungrafted nanoparticles. In Chapter 9 an anatase crystal size of 2.6 nm was

obtained at 250°C for 0.1 M TALD, but using only half the reactor volume and

with 0.15 M Sr2+ added to the reaction mixture the anatase crystal size was

reduced to 1.9 nm. The PXRD data (Figure 10.4A) shows the product is

maghemite and anatase. The 0.15 M Sr acetate elevated the pH and 0.1 M

acetic acid was added to reduce the pH to 5.

No crystalline SrTiO3 is formed due to the low pH, but from XRF a Sr/Ti

molar ratio of approximately 0.1 was obtained indicating that residual Sr is

present possibly on the titania surface. Barium acetate also gave a Ba/Ti

molar ratio of approximately 0.1. The single-heater-synthesis was conducted

using only the bottom heater such that the reaction mixture was quickly

quenched. The top heater was kept at room temperature, so as to recreate the

synthesis conditions of step 1.

If no base is added then SrTiO3/BaTiO3 should not form, and the anatase

crystals are simply further heated in the lower heater. Here the reaction

temperature was 360°C which was chosen to increase the reaction rate but

without the large decrease in density and shortened residence time in the

supercritical region. The exit temperature was measured by a thermocouple

to be a few degrees below 360°C. For this temperature profile using both

secondary heaters the anatase crystal size increased to 2.9 nm in the presence

of unreacted Sr(Ac)2 (Figure 10.4B). Switching to Ba(Ac)2 (Figure 10.4C) gave

an anatase crystal size of 2.2 nm.

These observations show that in the first reaction step anatase is indeed

formed using only a single heater, and that the addition of barium or

strontium salts does not change the anatase crystal size significantly. This is

important as an increase in anatase crystal size would slow down

SrTiO3/BaTiO3 formation. When NaOH is introduced the flow rate in the

lower reactor will increase and the residence time will drop to about 15


119

seconds which should cause the anatase crystal size to be less than 2.9 nm, if

no strontium or barium is present.

Figure 10.4 | PXRD of γ-Fe2O3@TiO2 synthesized at pH 5. A) Anatase crystal size is 1.9

nm using only a single heater in presence of unreacted Sr(Ac). B) Using two heaters (250°C

and 360°C, respectively) increases anatase crystal size to 2.9 nm with unreacted Sr(Ac)2. C)

Same as B but for Ba(Ac)2.

10.4.2 Step 2: Conversion of Titania to Sr- and Ba-Titanate

Ba(Ac)2

Figure 10.5A shows the PXRD data of the collected product when elevating the

pH by addition of 1.2 M NaOH from a separate pump to a γ-Fe2O3 and 0.1 M

TALD + 0.15 M Ba(Ac)2 + 0.1 M HAc mixture. It is interesting to observe that

the diffraction signal from the anatase crystal structure completely

disappears but that no BaTiO3 is formed. Only γ-Fe2O3 and BaCO3 are

present in the diffraction pattern. However, from XRF the Ti/Fe product

molar ratio is 1. Barium must therefore be incorporated into the titania

particles which disrupts the anatase crystal structure. This means that the

resulting BaxTiO2+x particles are amorphous, and that they must be grafted

on the iron oxides. If the titania nanoparticles were ungrafted and mobile


120

they would combine into larger crystals of BaTiO3. The absence of large

BaTiO3 nanocrystals suggests that TiO2 was successfully grafted on iron

oxide during step 1. STEM confirms the product as a nanocomposite (Figure

10.5B). The iron oxide core is successfully coated with a thin shell of a few

nanometers in agreement with the anatase crystal sizes.

Figure 10.5 | A) No crystalline BaTiO3 is formed using 0.1 M AIC, 0.1 M TALD, 0.15 M

Ba(Ac)2 and 0.1 M HAc, and diffraction peaks from anatase are absent. B) STEM-EDX of γ-

Fe2O3@BaxTiO2+x. Ba and Ti EDX signals are superpositioned, and Ba cannot be confirmed

in the shell.

EDX measurements are however unable to prove the shell composition

because the Ba L-lines (5-6 keV) overlaps with the Ti K-lines (5 keV). EDX

pictures of Ba and Ti would therefore always be completely identical even if

only one element was present. BaTiO3, BaCO3 and TiO2 can therefore not be

distinguished by EDX. From the TALOS software Ti is colored green and Ba

is coloured dark blue, but because it is the same data their superposition

(Figure 10.5B) is a light blue/turquoise. The colour legend was manually

changed to indicate Ti as this is the only element certain to be abundant in

the shell. From XRF the total Ba/Ti molar ratio is 1.1, and with the high

BaCO3 content the shell only contains minor quantities of barium.


121

Despite the strong indications from PXRD and XRF it is not possible to

confirm the core-shell nanoparticles as that of γ-Fe2O3@BaxTiO2+x by EDX.

The synthesis was therefore changed towards SrTiO3 which behaves

practically identical from a synthesis point of view.

Sr(Ac)2

The secondary precursor was switched from Ba to Sr, and the synthesis

parameters were changed slightly to further aid TiO2 to SrTiO3 conversion.

The Sr concentration was raised to 0.18 M Sr(Ac)2 while keeping the acetic

acid concentration fixed at 0.1 M. The NaOH concentration was raised to 2 M

to provide an even more alkaline environment for Sr incorporation into TiO2,

and the reaction temperature was elevated from 350°C to 360°C in the final

heater (Step 2).

The diffraction data (Figure 10.6A) shows the presence of γ-Fe2O3, SrCO3 and

SrTiO3. Figure 10.6B shows the PXRD pattern when washing the powder with

2 M HNO3. This is acidic enough for complete dissolution of SrCO3, but not

for γ-Fe2O3, TiO2 or SrTiO3. The SrCO3 phase is completely removed, while

the γ-Fe2O3 and SrTiO3 phases remain. Removing SrCO3 greatly improves the

data quality as its high content and scattering power otherwise dominates the

diffraction pattern.

Bubbles of CO2 were observed during acid treatment as the carbonates were

dissolved, which was completed after a few seconds. The treatment was

repeated and for extended dissolution time, but after quick centrifugation a

white suspension was visible above a grey precipitate. This was due to

smaller, SrTiO3 nanoparticles formerly grafted on the iron oxides which were

loosened by the acid treatment. Acid treatment was therefore avoided for

further investigations. A less concentrated acid may dissolve the SrCO3 and

not damage the hybrids, but this was not pursued further. The product can be

separated with a magnet which would reduce although not eliminate the


122

SrCO3 content, as the large SrCO3 particles are dragged along with the

magnetic hybrids.

From PXRD the average SrTiO3 crystal size was found to increase from 38

nm to 50 nm with HNO3 treatment which confirmed the loss of smaller

crystals. These crystal sizes are nevertheless much too large for a coating

thickness, and it is unlikely the maghemite particles should sit inside large

SrTiO3 single crystals.

Figure 10.6 | A) PXRD of γ-Fe2O3, SrTiO3 and SrCO3 synthesized at pH 5 followed by 2 M

NaOH. B) SrCO3 is completely removed after the acid treatment, but the SrTiO3 crystal size

increases.

STEM (Figure 10.7A) of the acid-treated product confirms that the large

SrTiO3 particles are ungrafted. The iron oxides are poorly coated, and Sr

forms ungrafted SrTiO3 from the mobile and ungrafted TiO2 particles faster

than from the fixed TiO2 particles (Figure 10.7B). Formation of ungrafted

particles is unexpected based on the selected synthesis parameters and may

be attributed to the acid treatment. Acetate is a surfactant, but it seems

unlikely that this minor increase in acetate concentration should induce a

surfactant effect that would shield the iron oxides from Ti4+ ions. A

10.5 Precipitation of carbonates

123

considerable factor in this synthesis is carbonate precipitation and its

formation is the subject of the next section.

Figure 10.7 | A) TEM of HNO3 treated γ-Fe2O3 and SrTiO3 from Fe:Ti:Sr molar ratio of

1:1:1.8. Large ungrafted cubic particles are observed. B) STEM-EDX of HNO3 treated γ-

Fe2O3@TiO2 for Fe:Ti:Sr molar ratio of 1:1:1.8. Titania is poorly grafted on iron oxides. Only

trace amounts of strontium is incorporated into the titania nanoparticles.

10.5 Precipitation of carbonates

A general problem when synthesizing BaTiO3 and SrTiO3 is the formation of

carbonates which precipitate above pH 6.131 Carbonates can be formed from

absorption of CO2 from the air, but in this synthesis the dominant carbon

source is citric acid from the Ammonium Iron Citrate (AIC) precursor which

decomposes into several organic species. The carbonate contribution from

TALD is insignificant. The NaOH concentration pumped at 10 mL/min was

increased gradually, and the crystalline TiO2 fraction was seen to decrease

with the emergence of SrCO3. This was not by titania removal as the Fe-Ti-Sr

product composition from XRF was comparable to the precursor

concentrations with the exception of 0.1 M and 0.2 M NaOH. Here the Sr/Ti

molar ratio was lower for the product than for the precursor which suggests

that parts of the Sr precursor was unreacted. This was confirmed by Rietveld


124

analysis which showed a much lower fraction of SrCO3 over iron oxide at low

NaOH concentrations.

The relative crystalline weight fractions obtained from Rietveld analysis are

plotted in Figure 10.8. A clear trend of decreasing anatase content with pH is

observed. SrCO3 is present as large ungrafted particles, and for the 1.0 M

NaOH synthesis a small fraction may have sedimented upstream of the outlet.

The main peaks from SrCO3 and anatase overlap and the anatase content is

likely overestimated at high pH. The data suggests that a concentration of at

least 1 M NaOH is required to disrupt the crystal structure of all the anatase

crystals. Higher concentrations seem unnecessary and may promote

precipitation of SrCO3.

Figure 10.8 | Relative weight fractions of crystalline SrCO3 and TiO2 relative to γ-Fe2O3.

Sr forms SrCO3 and converts anatase into amorphous SrxTiO2+x at alkaline conditions.

Product pH was 7, 11 and 14 for NaOH concentrations of 0.2 M, 0.5 M and 1 M,

respectively.

10.6 Evaluation of synthesis strategy

125

10.6 Evaluation of synthesis strategy

With these observations it is clear that SrCO3 precipitation is faster than Sr

diffusion into titania particles. The presence of SrTiO3 crystals is sporadic,

and after SrCO3 formation little Sr is available for incorporation.

The answer was to increase the strontium to carbonate ratio. One route is the

synthesis of maghemite (or a different magnetic core material) from an

inorganic precursor which would completely eliminate carbon. Despite its

appeal this route requires co-precipitation with base and titania would likely

nucleate homogeneously in the alkaline environment much like the mixed

SnO2 and TiO2 particles observed at high pH. Further pH modifications are

difficult and heterogeneous titania deposition is likely to suffer by this

method.

The easiest route was to simply decrease the AIC concentration. Moreover, it

was suggested in Chapter 9 that the anatase shell thickness could be

increased with TALD concentration. A thicker and denser anatase coating

would mean that the highest local anatase concentration is found in the shell

which should promote the formation of crystalline SrTiO3 grafted on the iron

oxide. A synthesis of 0.01 M AIC, 0.025 M TALD, 0.2 M Sr(Ac), 0.1 M HAc

and 2.0 M NaOH was performed. The selected precursor composition should

produce a thick anatase shell (Ti:Fe ratio of 2.5) with a large surplus of

strontium after SrCO3 precipitation. The lower heater temperature was set to

350°C.

Rietveld treatment of the PXRD data (Figure 10.9) shows the presence of γ-

Fe2O3 (19.7 nm), SrTiO3 (12.1 nm) and SrCO3. TEM (Figure 10.10A) shows the

product as isotropic particles with smaller cubic nanoparticles. Most of the

product consists of SrCO3, and an area free from large SrCO3 particles was

deliberately chosen as the high energy electron beam burns the SrCO3 during

which CO2 gas is released, which is undesirable for a vacuum technique. The

cubic particles form a close contact with the round particles and are not

preferentially agglomerated with other cubes. No ungrafted cubes are


126

observed. From TEM and STEM-EDX (Figure 10.10) the cubes are identified

as SrTiO3 particles attached to the iron oxide as if crystallized directly on

them. Individual EDX maps are displayed in Appendix 8.

Figure 10.9 | PXRD of γ-Fe2O3@SrTiO3 nanocomposite and SrCO3. The SrCO3 can be

removed by acid treatment, but this may destroy the composite structure.

The discrepancy between anatase and SrTiO3 crystal sizes indicate that

SrTiO3 formation is not formed by direct transformation of individual anatase

nanoparticles. A dissolution-precipitation model would promote formation of

large particles and not small SrTiO3 crystals of a narrow size distribution as

indicated by Figure 10.10A Different growth models have been proposed to

explain SrTiO3 formation from TiO2 nanoparticles, and different factors like

crystal structure, composition, particle size and morphology have been shown

to influence the reaction pathway.132, 133

Regardless of growth model it seems fair to assume that SrTiO3 is formed

where the local anatase concentration is highest. Such sites are found in the

shell, and SrTiO3 formation is likely initiated in the thickest parts of the shell.

If any anatase particles are ungrafted they are rapidly swept up by the

growing SrTiO3 particles. In the case of a fully closed titania shell it is

10.7 Effect of pH

127

possible that neighboring anatase particles are also swept up in a radius of a

few nanometers from the initial site of SrTiO3 formation. It seems plausible

then that there is a critical thickness below which crystalline SrTiO3 cannot

form in the shell.

Figure 10.10 | A) TEM of γ-Fe2O3@SrTiO3. Isotropic nanoparticles of 20-30 nm in size in

contact with smaller cubic nanoparticles. The cubic nanoparticles correspond in size with

the crystal size of SrTiO3. B) STEM-EDX of γ-Fe2O3@SrTiO3. The cubic SrTiO3

nanoparticles are in close contact with the iron oxides indicating that attractive forces exist

between the two oxides.

10.7 Effect of pH

The hybrid material in Figure 10.10 is not a core-shell particle, but it has a

large γ-Fe2O3/SrTiO3 interface for magnetoelectric interactions. Throughout

the chapter acids were used to neutralize the Ba and Sr salts for a targeted

pH of approximately 5 (the γ-Fe2O3@TiO2 hybrids in Chapter 9 where

synthesized at pH 6). At pH 5-6 the surface of the maghemite core is

electrostatically neutral (IEP of γ-Fe2O3 3.5-6.5) The observed morphology in

Figure 10.10 is therefore in accordance with the deposition mechanism based

on heterogeneous nucleation as determined from thermodynamics and

illustrated in Figure 8.1C.


128

The influence of pH was investigated by not adding acid to the secondary

precursor. Synthesis was performed using 0.01 M AIC, 0.03 M TALD, 0.2 M

Sr(Ac) and 1.4 M NaOH. The reaction temperatures of step 1 and 2 were

lowered to 230°C and 330°C, and the TALD concentration was increased to

promote a more uniform shell layer if possible. An initial reference was

synthesized in the absence of NaOH to measure the pH of the collected

product which was near pH 7.

From Figure 10.11 the product is a nanocomposite, and the morphology is a

core-shell structure. STEM-EDX confirms the presence of both Ti and Sr in

near equimolar concentrations in the shell. From the precursor

concentrations it is, however, evident that most of the added Ti and Sr species

Figure 10.11 | A-C) STEM-EDX maps of Fe, Ti and Sr. D) Superposition of A-C and image

from HAADF (not shown). A thin shell of SrTiO3 surrounds the magnetic iron oxide core.

10.8 Sub-conclusion

129

are not incorporated into the shell, but exists as ungrafted particles of SrTiO3

and SrCO3 which was also observed from STEM (Appendix 9). The pH of the

γ-Fe2O3TiO2 + Sr2+(aq) reference was neutral which is near the upper limit of

the maghemite IEP. The core particle surface was weakly negatively charged,

and the thin anatase shell fits with an electrostatic deposition mechanism. It

is unlikely that the TALD precursor is positively charged at neutral pH, but

transport of the molecule towards the negatively charged surface may be

facilitated by the ammonium ions or other positively charged group of the

TALD molecule.

10.8 Sub-conclusion

These proof-of-concept syntheses demonstrate that a multiferroic compound

can by synthesized hydrothermally in the dual-stage flow reactor, and that

the morphology is dependent of the pH. A nanocomposite of γ-Fe2O3 decorated

with SrTiO3 nanoparticles was successfully prepared when titania was

deposited at pH 5. At neutral pH the morphology changed to core-shell which

is likely due to changes in the electrostatic attractions between Ti(IV) species

and the negatively charged maghemite surface.

As a recipe for multiferroic nanoparticles the synthesis route is not complete,

and further studies are necessary to elucidate the titania coating mechanism

in the presence and absence of Ba/Sr salts. The main problem with the

pathway is the formation of large quantities of SrCO3. It is possible that TiO2

may still be deposited on γ-Fe2O3 at elevated pH using NaOH or NH4OH to

co-precipitate the iron oxide from an inorganic iron precursor which would

remove carbonates from the system. NH4OH is a weaker base than NaOH

and is likely the better candidate for a final pH that is not too high. For a

dedicated production line of this compound the synthesis would likely benefit

from splitting step 1 and 2 up into separate reactor stages and simply add the

Sr/Ba salts between titania formation and NaOH addition, but this requires

an additional cooler.


130

From the maghemite-free syntheses it was also shown that SrTiO3 and

BaTiO3 nanoparticles can be synthesized by addition of base to the newly

formed anatase suspension containing unreacted Ba/Sr salts. This allows

better tuning of the initial anatase crystal size and morphology which is

known to influence the formation of SrTiO3 and BaTiO3.

131

Chapter 11

MoS2

This chapter demonstrates the hydrothermal synthesis of molybdenum

disulphide. Sulphides are uncommon products from hydrothermal synthesis

because the reaction conditions generally favor oxides. MoS2 is an important

semiconductor with high catalytic performance in applications.

Phase pure MoS2 is first synthesized in a single-stage setup, and the crystal

size is shown to be independent of reaction temperature. However, the

material is not inert to higher temperatures, and it is demonstrated how

sulphur is increasingly removed from the particle surface.

MoS2 forms extended nanosheets, and it is shown how these can be used as a

support material for TiO2 nanoparticles by using the dual-stage configuration.

The best performances are obtained when the active particles are distributed

evenly over the substrate and neither material forms agglomerates with itself.

This chapter shows how the synthesis sequence can be used to prevent

agglomeration and form a dispersed product.

Chapter 11 MoS2

132

11.1 Structure and applications

MoS2 provides a rare example of a metal sulphide which can be readily

synthesized hydrothermally. It is an intensively studied semiconductor with

strong catalytic properties and of special interest to the petroleum industry in

which it is used as a hydrodesulphurization catalyst. It is furthermore used

for hydrogen evolution, chemical sensing and as electrode material for Li-

batteries due to its intercalation property. Traditional synthesis methods

include sulphidation of metallic Mo with H2S and vacuum techniques.134 Shi

et al.135 synthesized MoS2 from (NH4)2MoS4 by chemical vapour deposition.

Benck et al.136 synthesized amorphous MoS3 from Mo7O24(NH4)6·4H2O, H2SO4

and Na2S at room temperature. Dunne et al.137 used a hydrothermal flow

reactor to synthesize MoS3 from Mo7O24(NH4)6·4H2O, acetic acid and a heated

thiourea solvent stream, but extension of the reactor zone for increased

residence time was necessary for conversion of amorphous MoS3 into

crystalline MoS2.

By its layered structure it is an inorganic analogue to graphite and has been

used as a high-temperature lubricant in the past. It is interesting as a

support material, and experiments were carried out to graft anatase

nanoparticles on a MoS2 support during a MoS2&TiO2 synthesis where the

MoS2 supports were synthesized in the primary reactor and the TiO2

catalysts in the second reactor. For applications it is important that the

catalytically active material does not agglomerate into larger clusters but is

well-dispersed and grafted on the substrate. A TiO2&MoS2 synthesis was

therefore performed to investigate how the sequence of formation influenced

product characteristics including dispersibility.

11.2 Hydrothermal synthesis of phase pure MoS2

Synthesis of molybdenum disulphide required only the primary reactor stage

and the PRV was therefore installed just beneath the cooler of the primary

reactor. A 0.04 M aqueous solution of Ammonium thiomolybdate, (NH4)2MoS4


133

(ATM), was pumped at 5 mL/min and mixed with heated water at 15 mL/min

from the solvent HPLC pump.

For safety concerns an initial synthesis at 225°C was performed to test

whether high quantities of H2S gas and/or sulphuric acid were produced.

These by-products would compromise cooler performance and destroy the

reactor tube, respectively. Minor quantities of H2S was detected at the reactor

exit, but the gas was completely removed by placement of a fume hood above

the product collection bottle. No spurts of gas were observed, and the product

was pH neutral for all investigated temperatures. With these observations

the synthesis was deemed safe and the experiment continued. Reactor

residence time was 50-60 s.

Syntheses temperatures of 300°C, 350°C and 400°C all yielded hexagonal

structured MoS2 (Figure 11.1). A small hematite impurity is observed for the

synthesis at 400°C. This is not from reactor dissolution because pH was

neutral and no significant quantities of sulphuric acid could have been

produced. The iron oxide was likely produced during a previous synthesis and

is knocked loose from the reactor wall by the supercritical fluid. All

suspensions were dark, and the dried powder was a dark grey like that of

silicon.

Rietveld refinements showed large discrepancies between observed and

calculated intensities indicating that the particles are not spherical but

highly anisotropic. Platelets are a common motif for particles of hexagonal

crystal structure, and this morphology was confirmed by TEM analysis

(Figure 11.2A).

No significant change in diffraction profile was observed for the investigated

temperatures except for an increase in intensities. This is related to an

increase in diffracting volume. A highly crystalline powder will produce a

stronger diffraction signal compared to a less crystalline powder, but the

intensity also depends on how much material is being measured, and how it

is distributed across the irradiated sample holder. Although the crystallinity

Chapter 11 MoS2

134

is expected to increase with temperature the observed intensity increases are

not proof of this. For accurate crystallinity measurements a crystalline

internal standard must to be mixed with the powder like the procedure

employed for crystallinity measurements of TiO2 using Si as an internal

standard in 5.1.

Figure 11.1 | PXRD of MoS2 synthesized at 300°C, 350°C and 400°C. At 400°C a small

hematite impurity is observed, which is probably washed out from the reactor by the

supercritical fluid.

To account for the platelet morphology the diffraction data was modelled with

an anisotropic size model. The nanosheets were approximately 4.2 nm in the

crystallographic ab-plane and 2.4 nm along the 6-fold rotation axis (Table

11.1). These dimensions were constant for all investigated temperatures.

Figure 11.2B shows a TEM image of the sample synthesized at 350°C, and the

platelet size is at least an order of magnitude larger than the calculated

crystal size. This means that long-range order exists only locally within the

particle, and that the large nanosheets are polycrystalline. This has also been

reported in previous studies.137, 138


135

HR-TEM (Figure 11.3A) shows the lamellar structure of the nanosheets. The

sheets make folds at the edges, and the extra material increases electron

beam absorption making the edges appear darker. STEM-EDX (Figure 11.3B)

confirms that the particles are composed of Mo and S.

Crystalline MoS2 does not form immediately in the mixing point but only

after extended reaction time.137 It is therefore unlikely that crystal growth

should be arrested by neighboring crystals coming into contact and blocking

each other, because this growth mechanism would create a broad crystallite

size distribution which structural analysis does not support. The data

suggests that crystal growth is limited by an upper critical size based on the

constant crystallite size with temperature. From in-situ studies (NH4)2MOS4

decomposes into amorphous MoS3 at 150ºC, while crystalline MoS2 peaks are

not observed until at above 300ºC, however, the MoS2 crystallinity remain low

due to stacking faults and positional disorder.138

Figure 11.2 | A) TEM of MoS2 at 350°C. Particles form large agglomerates. B) TEM of

agglomerated MoS2 particles at 350°C. The platelet edges are darker due to curvature of

the platelets.

Chapter 11 MoS2

136

Figure 11.3 | A) HR-TEM of MoS2 at 350°C. The lamellar structure is visible by the large

interplanar lattice fringes. B) STEM-EDX of MoS2 at 350°C. Yellow signal indicate

presence of both S (green) and Mo (red).

MoS2 crystals originate from amorphous MoS3, but the transformations are

only a local event, and the result is polycrystalline sheets. The local S:Mo

composition in the regions separating the individual crystallites is likely

intermediates of MoS2 and MoS3. This is supported by XRF measurements

(Table ) which show a decrease in S:Mo molar content with temperature as

more and more MoS2 crystals form. This supports the observed increase in

crystallinity.

It is interesting to note that for the supercritical synthesis at 400°C the S:Mo

molar ratio is less than 2. A similar result was obtained for MoS2@TiO2 in

which MoS2 was first formed at 350°C and then reheated in the secondary

reactor for TiO2 formation. The added residence time produced a S:Mo molar

ratio of 1.5 comparable to that of the 400°C single-stage synthesis. It is

evident that sulphur is lost from the product with increasing temperatures.

This is either by substitution with oxygen in the strongly oxidizing

environment or by reduction of MoS2 to Mo2S3 or MoS.

By extrapolation it is likely that a S:Mo molar ratio of 2 exist around the

critical point of water. However, such a product would likely be a mixture of

11.3 TiO2 particles on MoS2 substrates

137

amorphous MoS3, crystalline MoS2 and sulphur deficient MoS2-x. Oxygen is

near the detection limit for the EDX detector and quantification is not

reliable because the signal overlaps with other light elements including

carbon from the film. Better oxygen estimates could be obtained by other

techniques such as X-ray Photoemission Spectroscopy.

Table 11.1 | Crystal sizes obtained from Rietveld analysis.

S:Mo molar ratio calculated from XRF spectroscopy

Crystal size (nm) S:Mo

(molar ratio) (a,b) plane c-axis

300°C 4.47(10) 2.43(10) 2.82

350°C 4.18(10) 2.38(10) 2.31

400°C 4.18(10) 2.37(10) 1.60

The product collected at 225°C could not be sedimented by centrifugation. It

was therefore evaporated and consisted of a dark brown precipitate with

small yellow islands. PXRD data (Appendix 10) revealed micron-sized

sulphur crystals (S8, alpha phase) and at least two large bumps at lower

diffraction angles meaning that the molybdenum sulphides are amorphous.

The dark brown colour is indicative of MoS3, and the low angle peaks are

likely related to the MoS3 interplanar spacing and unreacted precursor.


From the previous section it was demonstrated that nanosheets of

molybdenum disulphide can be synthesized hydrothermally, and that the

MoS2 crystal size is independent of reaction temperature. These observations

form the basis for the following investigation into MoS2-based nanocomposites.

For these syntheses the experimental setup was changed from single-stage to

standard dual-stage configuration (Figure 7.1).

Chapter 11 MoS2

138

Deposition of a catalytically active material on a substrate should a priori be

performed by first synthesizing the substrate and the catalyst second. MoS2

sheets were first synthesized in the primary reactor at 350°C as described in

section 11.2. A concentration of 0.04 M (NH4)2MoS4 was used in the primary

stage, and a corresponding concentration of 0.04 M TALD was used in the

secondary stage at 350°C (upper heater at 250°C). The TALD precursor was

mixed with acetic acid for a final concentration of 4 mM to lower the synthesis

pH, because the IEP of MoS2 is 1.9.139 Nonetheless, the collected product was

neutral, and the acid likely had no effect.

The sequence was afterwards reversed by swapping the two precursors; TiO2

was then synthesized in the primary reactor and MoS2 in the secondary

reactor. Flow rates and temperatures were identical for both syntheses. No

reference syntheses were made because no significant differences in crystal

sizes were expected.

PXRD data of MoS2&TiO2 and TiO2&MoS2 composites are shown in Figure

11.4. Both composites are composed of crystalline MoS2 and anatase. The

anatase crystal size was 5-6 nm in both compounds which is comparable to

the anatase crystal sizes from section 6.6 (TALD + Sn(NO3)4). Likewise, the

MoS2 crystal sizes were similar to the values in Table of single-phase MoS2.

From Figure 11.4 it is apparent, however, that large differences exist between

the intensities. In both cases the material formed in the primary reactor

forms the strongest diffraction signal. These differences in conversion

efficiencies are discussed later in the chapter.

From Figure 11.5 it is observed that TiO2 is not grafted on the MoS2 sheets in

the MoS2@TiO2 synthesis. The two materials form separate agglomerates and

are not evenly dispersed. The nanosheets are agglomerated shortly after

formation, and the TALD precursor crystallizes outside of the MoS2

agglomerates. By contrast, the reverse product, TiO2&MoS2, displays a much

different behavior towards agglomeration


139

Figure 11.4 | PXRD of composites of MoS2 and anatase in different synthesis sequence.

Higher crystallinities are obtained when the material is synthesized in the primary reactor.

Figure 11.5 | A) TEM of MoS2&TiO2. The two materials form separate agglomerates and

the material is not a composite. B) STEM-EDX of MoS2 rich region. Only a few anatase

particles can be seen. C) STEM-EDX of titania cluster. No MoS2 sheets are close by.

TEM (Figure 11.6) and STEM-EDX (Figure 11.7) of TiO2&MoS2 show the

platelets well dispersed among the titania particles. In fact, all the MoS2

sheets are covered by anatase particles while excess TiO2 particles form

agglomerates with other TiO2 particles. For this reason the support material

is always in volumetric excess when used in applications, but that is not the

Chapter 11 MoS2

140

purpose of these investigations. Supports are often made from light elements

(e.g. γ-Al2O3, carbon etc.), but the heavy molybdenum disulphide in excess

would completely dominate the PXRD data, and the lighter TiO2 would

therefore be difficult to characterize by PXRD.

It is interesting to note that no anatase particles are bonded to the MoS2 in

the case of the MoS2&TiO2 synthesis. If the anatase particles were attracted

to the sheets by van der Waal forces then this should be the case for both

samples. The data therefore suggests some other interaction is responsible for

the TiO2&MoS2 hybrid structure which likely occurs during MoS3

precipitation.

Figure 11.6 | TEM of TiO2&MoS2. Formation of MoS2 sheets after TiO2 is synthesized

yields a product of well-dispersed particles. No ungrafted MoS2 sheets are seen but

ungrafted TiO2 particles are present.


141

Figure 11.7 | STEM (A) and STEM-EDX (B) show that the TiO2 particles are closely

packed around the MoS2 sheets.

The tendency towards agglomeration or dispersion can be related to

observations made during the single-stage synthesis of MoS2 in section 11.2.

During these syntheses the exit stream would turn increasingly transparent

due to extensive particle agglomeration. These agglomerates worked like a

filter and caused a particle build-up inside the PRV and further upstream.

The PRV was therefore repeatedly opened and closed every minute or so to

create sudden fluctuations in pressure and flow speed to allow these large

agglomerates to exit. This problem became worse with temperature, and for

the final synthesis at 400°C the exit stream consisted of large macroscopic

black particles in a completely transparent supernatant.

Contrarily, the TiO2@MoS2 sample formed a very stable dispersion. Indeed, as

a side-note it should be mentioned that its sample for STEM analysis

(prepared from dried powder, dispersed in absolute ethanol and

ultrasonicated) is still fully stable at the time of writing, months after its

preparation. The MoS2@TiO2 STEM sample sedimented after a week, while

the 350°C single-phase MoS2 STEM sample (Figure 11.2 and Figure 11.3)

sedimented after a few days as was the case for most of the other STEM

samples.

Chapter 11 MoS2

142

As described earlier, large differences in PXRD intensities are observed in

Figure 11.1. Higher crystallinities of the primary material are always expected

due to the faster heating, higher average temperature and longer residence

time in the primary reactor, but this alone cannot account for such large

discrepancies. Ti and Mo precursors were added in a 1:1 molar ratio for both

syntheses, but from TEM images (Figure 11.5A and Figure 11.6) much of the

second phase appears to be missing. This is confirmed by XRF. The

MoS2&TiO2 has a Ti:Mo molar ratio of only 0.2 while the TiO2&MoS2 product

has a Ti/Mo ratio of 23. Poor MoS2 conversion is due to the short residence

time and slow heating, but the TiO2 yield is surprisingly low when

synthesized in a MoS2 suspension.

11.4 Sub-conclusion

Research into metal sulphides is greatly eclipsed by that into oxides. They are

difficult to produce on the nanoscale but hold many intriguing properties and

uses in applications. The synthesis of MoS2 under hydrothermal conditions is

therefore of great significance. Additionally, the used precursor is available

for other metals such as (NH4)2WS4, and W forms WS2 which is isostructural

with MoS2. It was discovered that the crystal size of MoS2 is constant with

reaction temperature, but that the S:Mo ratio decreases, and it is interesting

to see if the same effect applies for WS2.

The synthesis of MoS2&TiO2 clearly illustrated the issue of agglomeration.

Attempts of making a composite material from a simple mixture of the

premade powders or suspensions would likely lead to similar ungrafted and

agglomerated particles. This is indeed a major problem for supported

materials as inadequately dispersed catalysts are both expensive and

inefficient. The results in this chapter demonstrate that sequential synthesis

of the two materials can lead to a well-dispersed product, and that this may

be readily accomplished using the dual-stage flow reactor.

11.4 Sub-conclusion

143

Titania was most evenly distributed when synthesized first and further

investigations may look into replacing TALD with TTIP as the choice of

precursor have been shown to significantly change the formation chemistry

including the preference towards heterogeneous structures. The results open

up for a range of other interesting combinations for supported heterogeneous

catalysts for instance SiO2 or γ-Al2O3 as support for metallic nanoparticles

(e.g. platinum, silver, gold, copper) which have not been addressed in this

dissertation but are of great technological importance.

145

Chapter 12

Concluding Remarks

This dissertation demonstrates how materials can be combined into hybrids

on the nanoscale, and how these remarkable nanocomposites can be produced

by hydrothermal flow chemistry. By design, the dual-stage flow reactor

(Chapter 3) promotes the conditions responsible for heterogeneous deposition

by dividing formation of the core and the shell material into separate heating

zones. Synthesis of TiO2@SnO2 (Chapter 6) proves the advantages of this

design, and the main parameters responsible for formation of hybrid or mixed

nanoparticles are mapped out.

The suspension pH is of key importance. It directs the shell morphology into

thin, uniform and closed shells or thick, irregular and incomplete shells using

silica as an example (Chapter 7). The interplay of pH and shell morphology

led to the development of an electrostatically driven deposition model for

predicting synthesis conditions leading to the core-shell morphology (Chapter

8). Uniform core-shell nanoparticles can be synthesized by relating the

isoelectric point with the synthesis pH, but the reaction rate of the shell

precursor has to be accounted for as observed for α-Fe2O3@SiO2.

A shell material is formed from interactions of the shell precursor with the

core, and this may induce changes in both the crystal structure and size of the

shell crystallites (Chapter 6, 8). Directing the shell precursor into a targeted

crystalline phase can be hampered by these interactions and controlling

theshell composition requires special attention as demonstrated by the three

Chapter 12 Concluding Remarks

146

products γ-Fe2O3@CuxO, γ-Fe2O3@NiOxLy and γ-Fe2O3@Zn6(OH)6(CO3)2

(Chapter 8).

Furthermore, it was shown how anatase can be deposited on a magnetic

nanoparticle (Chapter 9). The electrostatic model demonstrated how thin

shells are formed from repulsion of shell and shell precursor, but that the

shell thickness can be increased if at least one component is electrostatically

neutral (Chapter 9, 10). An anatase shell was demonstrated as an

intermediate for γ-Fe2O3@SrTiO3 multiferroic nanocomposites (Chapter 10).

Different morphologies can be obtained; from SrTiO3 crystals grafted on the

magnetic core, to a thin and uniform SrTiO3 nanoshell surrounding the core,

but the synthesis pathway is complex. The choice of both core and shell

precursors were shown throughout the dissertation to be of great importance

for whether a hybrid or a mixed product is formed.

Finally, it was shown how MoS2 (Chapter 11) can be synthesized at

hydrothermal conditions. Platelets can be used as support material for

multiple smaller nanocatalysts, and the dissertation demonstrate how this

class of nanocomposites can also be synthesized in the dual-stage reactor

without agglomeration of the support.

The dual-stage flow reactor is a highly versatile reactor type that allows a

nanoparticle to be functionalized in high yields for use in applications. This

dissertation maps the main parameters responsible for nanocomposite

formation and describes how different synthesis conditions can direct the

shell morphology. In combination, the dual-stage reactor and this dissertation

serve to provide a green and sustainable route towards synthesis of these

highly advanced hybrid nanomaterials.

147

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Appendix

A1

Appendix

Additional data is enclosed in the following pages which supports the findings

of the dissertation.

Appendix

A2

Appendix 1 3D Rendering of Model of Dual-Stage Flow Reactor

Figure A1| 3D model of backside of dual-stage flow reactor.

Appendix

A3

Figure A2| 3D model of front side of dual-stage.

Appendix

A4

Appendix 2 TiO2 unit cell sizes

Anatase unit cell sizes of selected TiO2 syntheses. The size of the anatase unit

cell is unchanged for the TiO2@SnO2 nanocomposite indicating that Sn does

not substitute Ti. The larger unit cell size of Sn substituted anatase is shown

for comparison. A Sn substituted anatase compound, Ti0.85Sn0.15O2, is

included for comparison.

a / Å c / Å V / Å3

TiO2, 0.25 M, 5% H2O, 300ºC 3.7834(1) 9.5062(4) 136.072(9)

TiO2, 0.25 M, 100% H2O, 350ºC 3.7801(2) 9.4903(8) 135.61(2)

TiO2, 0.25 M, 10% H2O, 400ºC 3.7822(1) 9.4987(4) 135.882(8)

TiO2, 1.0 M, 10% H2O, 350ºC 3.7842(2) 9.5088(4) 136.167(1)

TiO2@SnO2 3.7873(6) 9.485(2) 136.05(4)

Ti0.85Sn0.15O2 (ICSD-72769) 3.8065(2) 9.6733(8) 140.16(2)

Appendix

A5

Appendix 3 ICP-OES of SnO2 and TiO2@SnO2 from SnCl4

Supernatants of SnO2 and TiO2@SnO2 products from the alkaline syntheses

were analyzed with ICP-OES. Large quantities of unreacted Sn were detected.

Sn conversion increases with temperature. The addition of titania does not

promote Sn precipitation at high pH.

Synthesis Ti

(mg/kg)

Sn

(mg/kg)

Ø@SnO2, alkaline, 200°C 0.009 > 42.919

Ø@SnO2, alkaline, 250°C 0.007 > 13.908

Ø@SnO2, alkaline, 300°C 0.001 2.001

TiO2@SnO2, alkaline, 300°C 0.016 > 3.345

Appendix

A6

Appendix 4 PXRD of Sn:Ti molar ratio of 10

The SnO2 dominates the diffracted signal due to the higher scattering power

and higher contents over titania.

Figure A3| PXRD of TiO2&SnO2 for a Sn:Ti molar ratio of 10.

Appendix

A7

Appendix 5 FTIR spectra of SiO2

Figure A4| FTIR spectra of γ-Fe2O3@SiO2

and γ-Fe2O3. The transmittance is lowered

because the powder is black.

Figure A5| FTIR spectra of TiO2@SiO2 and

TiO2.

Figure A6| FTIR spectra of α-Fe2O3@SiO2 and α-Fe2O3.

Appendix

A8

Appendix 6 α-Fe2O3@SiO2: Temperature and Concentration

Figure A7| TEM of α-Fe2O3@SiO2 at 300°C from 0.1 M Fe(NO3)3 and 0.2 M TEOS. Shell

morphology is similar to that of the product obtained at 100°C.

Figure A8| TEM of α-Fe2O3@SiO2 at 300°C from 0.1 M Fe(NO3)3 and 1.0 M TEOS. Iron

oxide nanoparticles are incorporated into micron-sized aggregates of silica due to the high

Si:Fe molar ratio.

Appendix

A9

Appendix 7 PXRD of γ-Fe2O3@TiO2, 350°C

The higher reaction temperature of the synthesis at 350°C increases the

anatase crystal size 25% compared to the product obtained at 250°C. The high

secondary reaction temperature causes the maghemite crystal size to increase

slightly from 17.6 nm (250°C ) to 18.1 nm (350°C)

Figure A9| PXRD of γ-Fe2O3@TiO2 at 350°C and Ti:Fe molar ratio of 1.

Appendix

A10

Appendix 8 STEM of γ-Fe2O3@SrTiO3

Figure A10| STEM-EDX of γ-Fe2O3@SrTiO3.

Appendix

A11

Appendix 9 γ-Fe2O3@SrTiO3 without acid for pH control

Figure A11| STEM-EDX of γ-Fe2O3@SrTiO3. Most of the SrTiO3 particles are ungrafted as

~10 nm large particles. A thin shell of SrTiO3 surrounds the magnetic particles. The product

also contains SrCO3 (not shown)

Appendix

A12

Appendix 10 PXRD of (NH4)2MoS4 225°C decomposition product

Figure A12| PXRD of products obtained at 225°C for hydrothermal flow synthesis of

molybdenum sulphides. Sharp peaks belong to sulphur but a few peaks do not correspond to

the alpha phase. Large bumps at lower angles most likely correspond to amorphous MoS3

interplanar distances and unreacted precursor.

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