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2017 Origins of Unconventional Magnetism in Coinage Metal Nanomaterials

Marenco, Armando J.

Marenco, A. J. (2017). Origins of Unconventional Magnetism in Coinage Metal Nanomaterials (Unpublished doctoral thesis). University of Calgary, Calgary, AB. doi:10.11575/PRISM/27233 http://hdl.handle.net/11023/3610 doctoral thesis

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Origins of Unconventional Magnetism

in Coinage Metal Nanomaterials

by

Armando J. Marenco

A DISSERTATION

SUBMITTED TO THE FACULTY OF GRADUATE STUDIES

IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE

DEGREE OF DOCTOR OF PHILOSOPHY

GRADUATE PROGRAM IN CHEMISTRY

CALGARY, ALBERTA

January, 2017

c Armando J. Marenco 2017 Abstract

The studies presented in the dissertation focused on the unconventional magnetic properties of coinage metals - Cu, Ag, and Au - nanomaterials synthesized in the gas phase by sputtering. Unlike other common synthetic methods, gas-phase synthe- sis creates nanoparticles and thin films free of capping ligands allowing for pristine surface studies. The two types of nanomaterials synthesized were sub-12 nm diame- ter nanoparticles and thin films. These three parameters, namely the nature of the coinage metal, unmodified surfaces, and nanodimensionality, were the core effects independently studied.

The unconventional magnetism of Cu nanomaterials has been studied in highly- pure and capping-ligand-free nanoparticles and thin films. Superconducting quantum interference device (SQUID) room-temperature (300 K) measurements displayed no size correlation to the ferromagnetic behavior observed in Cu nanoparticles ranging from 4.5 ± 1.0nmto9.0± 1.8 nm in diameter. Moreover, magnetic quartz crystal microbalance (MQCM) in situ tests of 4.5 ± 1.0 nm nanoparticles under vacuum con- ditions showed magnetic behavior only after the onset of oxidation. SQUID analysis conducted on Cu thin films exposed to several heat treatments demonstrated minor oxidation inducing higher ferromagnetic responses compared to extended oxidation.

i Further analysis of nanomaterial samples exhibiting the highest magnetic responses

indicated an atomic ratio of ∼3-5:1 Cu:O suggesting non-stoichiometric oxidation as the source of the ferromagnetic signature.

Similar ferromagnetic results were obtained for Ag nanomaterials. No size cor- relation to magnetism in Ag nanoparticles ranging from 3.3 ± 0.9nmto7.8± 1.3

nm was determined by SQUID magnetometry. Additionally, MQCM under vacuum

conditions of 3.3 ± 0.9 nm Ag NPs shows magnetic behaviour only after the onset

of oxidation. The trend in SQUID magnetometry shows a higher saturation mag-

netization for samples exposed to oxygen compared to inert atmospheres, which is

opposite to the Cu findings. However, this also indicates non-stoichiometric oxida-

tion at the surface as the reason for the observed magnetism as supported by an ∼8:1

Ag:O ratio from MQCM measurements. Finally, heat-treated Ag thin films present

a lower saturation magnetization compared to those kept in oxygen without heating.

This latter observation could be the result of driving the oxidation to stoichiometric

surface oxides AgO and Ag2O which are known to be diamagnetic. For Au nanomaterials, our findings show some promise but are inconclusive. A

single diameter was synthesized at 3.9 ± 1.7 nm. Unlike the MQCM studies observed

for Cu and Ag, Au NPs show no significant signal change which could be interpreted

as ferromagnetic behaviour. However, SQUID magnetometry does provide a clear

ferromagnetic signal. Further studies are required to determine for certain if oxidation

plays a role in Au nanomaterials as determined for Cu and Ag. Acknowledgements

First and foremost, I would like to thank my supervisors Prof. Simon Trudel and

Dr. David B. Pedersen for their guidance, encouragement, and support throughout this lengthy and successful process. Also, I would like to thank the members of my committee Prof. George K. H. Shimizu and Prof. Venkataraman Thangadurai whom provided me with excellent advise during the entirety of my program. Additionally, I am thankful to Prof. Colin Dalton and to Prof. Andrew P. Grosvenor for reviewing my thesis and providing helpful comments.

I am grateful to our collaborators Prof. Jillian M. Buriak at the University of

Alberta, Prof. Frank C. J. M. van Veggel at the University of Victoria, and to

Prof. Mark J. MacLachlan at the University of British Columbia. Our collaborations allowed me to greatly expand my knowledge on magnetism.

The acknowledgments could not be completed without thanking current and past members of the Trudel Group, specially to Raphael Dong, Jennifer Emara, Luvdeep

Bhandari, and Casey Platnich.

I am grateful to my parents, my siblings and their families for all their support.

Finally, I would like to thank the entire support personnel in the Department of the Chemistry at the University of Calgary.

iii Dedicated to my nieces and nephews. Contents

Abstract ...... i Acknowledgements ...... iii Contents...... v ListofTables...... viii ListofFigures...... ix ListofSymbols...... xi 1 Introduction...... 1 1.1Originofmagnetism...... 1 1.2Classificationofmagneticmaterials...... 4 1.2.1 Diamagnetism...... 4 1.2.2 Paramagnetism...... 7 1.2.3 Collective Magnetism ...... 9 1.3ExchangeInteractions...... 15 1.3.1 Exchange...... 15 1.3.2 DirectExchangeInteraction...... 18 1.3.3 IndirectExchangeInteractions...... 19 1.4 Conventional and Unconventional Magnetism ...... 22 1.4.1 Conventionalmagnetism...... 22 1.4.2 Unconventionalmagnetism...... 25 1.5 Unconventional Magnetism in Coinage Metals ...... 26 1.6Nanomaterials...... 31 1.6.1 Nanoparticles...... 32 1.6.2 Thinfilms...... 33 1.7Sputteringprocess...... 33 1.7.1 Sputteringtheory...... 34 1.7.2 Gas-phasenanoparticlegeneration...... 36 1.8 Research goal ...... 38 1.8.1 Importanceofmagnetism...... 41 1.9Outlineofthesis...... 42 2 Experimentaldetails...... 44 2.1Nanomaterialsynthesisinthegasphase...... 44 2.1.1 Nanoparticlepreparation...... 45

v 2.1.2 Thinfilmpreparation...... 51 2.2Opticalsetup...... 55 2.3 Magneto-optical Kerr effect setup ...... 55 2.4 Superconducting quantum interference device magnetometry . . . . . 57 2.5 Magnetic quartz crystal microbalance ...... 61 2.5.1 In situ MQCM...... 63 2.6X-rayphotoelectronspectroscopy...... 66 2.7Atomicforcemicroscopy...... 66 2.8Scanningelectronmicroscopy...... 66 2.9PowderX-raydiffraction...... 67 2.10 Experimental methodology for Cu nanomaterials ...... 67 2.11 Experimental methodology for Ag nanomaterials ...... 68 3 Opticalstudiesofsilvernanomaterials...... 71 3.1Introduction...... 71 3.2 Magneto-optical effects ...... 72 3.2.1 Faradayeffect...... 72 3.2.2 Kerreffect...... 75 3.3 Magneto-optical Kerr effect ...... 79 3.4 An alternative to direct measurement of the Kerr rotation ...... 82 3.5Resultsanddiscussion...... 83 3.5.1 Preliminaryopticalexperiments...... 83 3.5.2 MOKEmeasurements...... 90 3.6Conclusion...... 95 4 Inducing ferromagnetic behaviour in Cu nanomaterials ...... 96 4.1Introduction...... 96 4.2Experimental...... 99 4.3Results...... 99 4.3.1 CuNPs...... 99 4.3.2 Cufilms...... 102 4.3.3 Cu NPs - in situ MQCMmeasurements...... 103 4.4Discussion...... 115 4.4.1 Sizedependence...... 115 4.4.2 Oxidation...... 116 4.4.3 Our magnetic system compared to other systems ...... 116 4.5Conclusion...... 120 5 On the origin of ferromagnetic signature in Ag nanomaterials . . . . . 122 5.1Introduction...... 122 5.2Experimental...... 126 5.3Results...... 126 5.3.1 AgNPs...... 126 5.3.2 Agfilms...... 129 5.3.3 Ag NPs: In Situ OxidationEffects...... 131 5.3.4 SurfaceAnalysis...... 136 5.4Discussion...... 138 5.4.1 SurfaceSpeciesCharacterization...... 138

vi 5.4.2 OriginofferromagnetisminAgNPs...... 147 5.5Conclusion...... 150 6 Preliminarystudiesongold...... 151 6.1Introduction...... 151 6.2Experimental...... 153 6.2.1 SynthesisofAunanomaterials...... 153 6.2.2 Otherprocedures...... 154 6.3Resultsanddiscussion...... 154 6.3.1 AuNPsandthinfilms...... 154 6.3.2 MQCM...... 155 6.3.3 Superconducting quantum interference device magnetometry . 157 6.3.4 PowderX-raydiffraction...... 158 6.4Conclusions...... 159 7 Finalconclusionsandfuturedirection...... 161 7.1Summaryandfinalconclusions...... 161 7.2Futuredirection...... 163 7.2.1 X-ray magnetic circular dichroism ...... 163 7.2.2 In situ magneto-opticalKerreffect...... 164 7.2.3 Reductionofnanomaterials...... 164 7.2.4 Solid-state 109Ag nuclear magnetic resonance ...... 165 7.2.5 Magnetic signal longevity ...... 165 7.2.6 Surfacerearrangement...... 166 7.2.7 SharperNPdiameters...... 166 7.2.8 Othersyntheticmethods...... 166 Bibliography...... 168 A AppendixA...... 192

vii List of Tables

Table 2.1 NP generation conditions during testing of parameters . . . . . 49 Table2.2ConditionsforthegenerationofCuNPs...... 68 Table2.3ConditionsforthegenerationofAgNPs...... 69

Table 3.1 Kerr rotation θK literature values for Fe, Co and Ni ...... 83 Table 3.2 Conditions for the generation of 3.7 nm Ag NPs ...... 86

Table 4.1 XPS Cu 2p3/2 peakpositioning...... 114 Table 5.1 Magnetic properties of Ag nanomaterials ...... 136 Table 5.2 Ag 3d 5/2 XPSfittinganalysisforAgthinfilms...... 146

Table 6.1 Literature M s valuesforseveralAuNPs...... 152 Table6.2ConditionsforthegenerationofAuNPs...... 154 Table 6.3 Magnetic properties of Au NPs ...... 157

viii List of Figures

Figure 1.1 Scheme of an electron orbiting its nucleus ...... 2 Figure 1.2 Magnetic moments in diamagnetic and paramagnetic materials 6 Figure 1.3 Density of states for paramagnetic materials ...... 9 Figure 1.4 Magnetic moments in collective magnetism materials . . . . . 11 Figure 1.5 Typical Mvs.H loopsofseveralmaterials...... 13 Figure1.6TheBathe-Slatercurve...... 18 Figure 1.7 NiO as a superexchange interaction example ...... 20 Figure 1.8 Fe3O4 as a double exchange interaction example ...... 20 Figure1.9RKKYinteractionplot...... 22 Figure 1.10 Stoner model for 3d ferromagnetism...... 24

Figure2.1Gas-phasedepositionsystem...... 45 Figure 2.2 The aggregation zone during sputtering ...... 47 Figure2.3NPsizeeffectofdifferentparameters...... 50 Figure 2.4 Cu metal targets exhibiting nanostrusture deposition . . . . . 52 Figure 2.5 Scenarios depicting nanomaterial redeposition onto target . . . 53 Figure2.6Opticalsetupphotograph...... 56 Figure2.7MOKEsetupphotograph...... 57 Figure2.8MOKEsetupphotograph...... 58 Figure2.9SQUIDmajorcomponents...... 60 Figure2.10MQCMconcept...... 63 Figure 2.11 In situ MQCMsetup...... 65 Figure 2.12 Nanogenerator and collection locations for Cu nanomaterials . 69 Figure 2.13 Nanogenerator and collection locations for Ag nanomaterials . 70

Figure3.1TheFaradayeffect...... 73 Figure 3.2 Polarized light and Faraday rotation ...... 74 Figure3.3AbsorbedlightandKerrangle...... 75 Figure 3.4 Microscopic quantum source of the Kerr effect ...... 78 Figure3.5ExperimentalMOKEconfigurations...... 81 Figure3.6PolarMOKEsetup...... 81 Figure 3.7 Placement of NaCl prisms during thin film collection . . . . . 85

ix Figure 3.8 Set-up for preliminary optical experiments ...... 86 Figure3.9Massspectrumof3.7nmAgNPs...... 87 Figure 3.10 Absorbance spectra of s-andp-polarized light by Ag NPs . . 88 Figure 3.11 NPs interaction with s-andp-polarizedlight...... 89 Figure 3.12 Absorbance spectra of s-andp-polarized light by Ag thin films 90 Figure3.13Specularanddiffusereflections...... 91 Figure 3.14 MOKE measurements for Ag thin film and NPs ...... 94

Figure 4.1 Cu NPs diameters and SQUID measurements ...... 101 Figure 4.2 SQUID measurements of Cu NPs stored in Ar and O2 ..... 102 Figure 4.3 SQUID measurements of Cu thin films stored in Ar and O2 . . 104 Figure 4.4 SQUID measurements of heat-treated Cu thin films ...... 105 Figure 4.5 MQCM proximity effect and proof of concept ...... 106 Figure4.6RawMQCMdataofCuNPs...... 107 Figure 4.7 MQCM and saturation magnetization data of Cu NPs . . . . . 108 Figure4.8XPSspectrumsurveysofCuthinfilms...... 109 Figure 4.9 S 1s XPSregionofCuthinfilms...... 110 Figure 4.10 Cu 2p and O 1s XPS regions of Cu thin films stored in gas . . 111 Figure 4.11 Cu 2p and O 1s XPS regions of Cu heat-treated thin films . . 112 Figure4.12XPSdataofCuthinfilms...... 114 Figure 4.13 Unit cells of metallic Cu, Cu2OandCuO...... 117 Figure 4.14 Proposed oxidation stages of Cu ...... 120 Figure 4.15 Proposed mechanism for Cu oxide formation on nanomaterials 121

Figure 5.1 Ag NPs diameters and SQUID measurements ...... 127 Figure5.2GaussianfittingforAgNPdistributions...... 128 Figure 5.3 ON- and OFF-target Ag thin films M vs H comparison . . . . 130 Figure 5.4 SQUID measurement for a ON-target Ag film ...... 131 Figure 5.5 ON- and OFF-target Ag thin films SEM and AFM comparison 132 Figure 5.6 ON- and OFF-target Ag thin film SEM comparison ...... 133 Figure5.7MQCMdataofAgNPs...... 134 Figure 5.8 SQUID of Ag NPs and thin films stored in Ar and O2 ..... 135 Figure 5.9 SQUID measurements of heat-treated Ag thin films ...... 137 Figure5.10EDXSofAgthinfilm...... 138 Figure5.11XPSspectrumsurveysofAgthinfilm...... 139 Figure 5.12 Ag 3d and O 1s XPSregionsofAgthinfilms...... 140 Figure 5.13 Literature Ag 3d5/3 XPSpeakcomparison...... 143 Figure 5.14 Proposed oxidation stages of Ag ...... 149

Figure 6.1 Literature Ms valuesforAuNPs...... 153 Figure6.2MassspectrumofAuNPs...... 155 Figure6.3MQCMofAuNPs...... 156 Figure 6.4 SQUID of Au NPs and thin films stored in Ar and O2 ..... 158 Figure 6.5 XRD of ON-target Au thin films ...... 159

x List of Symbols, Abbreviations and Nomenclature

Symbol Definition

ΔfMQCM frequency change during MQCM θ angle, see context for definition

μ magnetic moment

μB Bohr magneton

νplasma plasma oscillation

ρ density

χ magnetic susceptibility

 reduced Planck’s constant or h/2π

AFM atomic force microscopy

E energy

e charge of the electron

EDXS energy dispersive X-ray spectroscopy

E ex exchange energy

emu electromagnetic unit of magnetic moment

H external magnetic field strength

H c coercivity

xi I current ICD inert gas condensation

K Kelvin M magnetization me mass of the electron MOKE magneto-optic Kerr effect

MQCM magnetic quartz crystal microbalance

M s, M r saturation, remnant magnetization

NA Avogadros number NP nanoparticle

PXRD powder X-ray diffraction

QCM quartz crystal microbalance

SAM self-assembled monolayer

sccm standard cubic centimeter per minute

SEM scanning electron microscopy

SOC spin-orbit coupling

SPR surface plasmon resonance

SQUID superconducting quantum interference device

THPC tetrakis(hydroxymethyl)phosphonium chloride

V crater volume of sputtered material XMCD X-ray magnetic circular dichroism

XPS X-ray photoelectron spectroscopy

Y sputtering yield

xii Chapter 1

Introduction

1.1. Origin of magnetism

In the strictest sense, magnetism originates from electrons – specifically from unpaired

electrons. Atomically, the spatial motion of an electron about the nucleus creates

a magnetic moment due to the electron’s negative charge (see Figure 1.1). The

classical mechanics equivalent of this phenomenon is having an electric or Amperian

current (i.e., a moving charge) circulating on a wire which creates a magnetic field

perpendicular to the wire. In Figure 1.1, the physical orbital angular momentum created by the rotational motion of the electron around the nucleus creates an upwards force, while the magnetic orbital angular momentum is created by any moving charged particle. By convention, the direction of the current is opposite to that of the electron which creates the downward orbital magnetic moment. Mathematically, the orbital

z angular magnetic moment μorbital is defined by the following equation

z −μB μorbital =  Lz (1.1)

where the superscript and subscript z refers to the quantization axis described by the

1 Figure 1.1: Scheme of an electron orbiting its nucleus. The counter clockwise orbiting motion (blue arrow) of the electron (red sphere) creates an upward angular momentum (Lz). By convention the current’s direction is opposite to the electron’s motion. The clockwise circulating current (yellow arrow) induces a downward orbital angular z momentum (μorbital). Adapted from Refs. 1 and 2.

magnetic field along this particular direction, μB is the Bohr magneton,  is Planck’s

constant divided by 2π,andLz is the orbital angular moment projected onto the

quantization axis. The Bohr magneton μB is defined as the magnetic moment of a free electron (i.e., an unbound electron will induce no orbital angular momentum)

and is defined as:1

 e −24 −1 μB = =9.27 × 10 J · T (1.2) 2me

where e is the charge of an electron and me is the mass of the electron.

z A second component to the total magnetic moment μtotal is the spin angular

momentum. This second component is an intrinsic property of the electron (similarly

to its mass and charge); and, it is not an actual spinning motion.3 The spin is a purely

quantum phenomenon with no classical analogue and should be taken as a fact4 as

it has been experimentally measured in the famous Stern-Gerlach experiment. The

z magnetic moment associated with the spin μspin is defined by :

z − μB μspin = ge  Sz (1.3)

5 where g e is the g-value for a free electron and by definition has a values of 2.0023, and

2 Sz is the spin angular moment projected onto the quantization axis. Other subatomic particles possessing spin are protons and neutrons; however, due to their larger mass in comparison to electrons,3 their spin contribution to the total magnetic moment

z μtotal in an atom is negligible and disregarded. For example, the magnetic moment of

1 nuclei is smaller by a factor of 1836 when compared to μB. As seen in Equation 1.2, the mass of subatomic particles is present at the denominator which would result in

a smaller value for μspin with increasing mass. This particular 1836 factor results

from the mass of the electron and proton being 9.109 × 10−31 and 1.673 × 10−27 kg,

respectively. Thus, the total magnetic moment of a material is provided by both the

orbital angular moment and the spin angular moment of electrons:

z z z −μB μtotal = μorbital + μspin =  (Lz + geSz) (1.4)

Moreover, the contributions to the total magnetic momenta from both the or-

bital momentum and from the spin momentum are not equal.2 Free atoms with

open shells carry spin (μspin) and orbital (μorbital) magnetic moments. However, as

atoms form bulk structures, both momenta are substantially attenuated.6 Further-

7 more, μorbital can be reduced to 5-10% of the total magnetic moment or can be

completely quenched6 in macroscopic systems by the symmetry of the crystal lattice

6,7 yielding an orbital angular momentum of zero, Lz =0. This effect is the result of

the orbital magnetic moment being influenced by orbital hybridization and electron

delocalization.6

Finally, magnetic moments are additive and the overall contribution of all the

atomic magnetic moments in a material defines its magnetic behaviour. For instance,

the overall contributions can be nil in cases of atoms containing fully occupied orbitals.

An atom exhibiting closed shells has zero orbital angular momentum because all

orbital angular momenta sum to zero3 as electrons orbiting in opposite direction1

3 cancel each other. Additionally, a closed shell indicates that all electrons are paired and there is no net spin3 due to pairing up with opposite spins (i.e., ↑↓) as stated by the Pauli exclusion principle.1 It is important to note that the total magnetic moment can be modeled vectorially;5 thus, both the orbital and spin angular momenta are considered vector quantities. As will be discussed in Section 1.3, when magnetic moments coupled with each other, they form magnetic interactions known as exchange interactions.

1.2. Classification of magnetic materials

As mentioned earlier, magnetism originates from the electron. However, within a bulk structure, electrons are not isolated and may sometimes interact with electrons residing on the same atom or adjacent atoms. The presence of electronic interaction gives rise to the physical phenomena of magnetism. Although the magnetic behaviour of any material can be described as either magnetic or non-magnetic simply based on their physical response to, for example, an external magnetic field, their classifi- cation is more complex due to the different electronic interactions present within the material.

1.2.1. Diamagnetism

Since magnetic moments are considered vector quantities, when all electrons in a material are oriented in such a fashion that they cancel one another, the atom as a whole has no net magnetic moment.4 A material composed entirely of these atoms is considered a diamagnetic material. Conversely, an atom consisting of paired and unpaired electrons exhibits only a partial cancellation of momenta (i.e., non-zero), in turn experiencing a net magnetic moment. Thus, diamagnetism is a property of all materials8 as all materials, whether exhibiting magnetism or not, contain atoms with

4 core electrons occupying full orbitals; however, diamagnetism is usually only evident when other magnetic effects are not present.1

In general, diamagnetism is purely an induction effect caused by Lenz’s law.8 As an external magnetic field H is applied to a diamagnetic substance, a magnetic flux is induced within it counteracting H, and in the process creating magnetic dipoles oriented antiparallel to H resulting in a negative magnetic susceptibility χ value (see

Figure 1.2a).8 The magnetic susceptibility χ is the ability of a material to change its magnetization M with a changing applied external magnetic field H and is defined as:9

M χ = (1.5) H The concept of diamagnetism can be explained with a two electron system,1 in which an external magnetic field H exerts a magnetic force on the orbiting electrons which already are experiencing an electrostatic attraction from the nucleus. This new magnetic force causes the electron orbiting antiparallel to the magnetic field to speed up, while decreasing the speed of the electron orbiting parallel to it.1 (Note that in the absence of the external magnetic field, the orbiting speed of both electrons is the same but in opposite direction resulting in a cancellation of the orbital angular magnetic moment μorbital.) The unequal speeds results in a net magnetic moment opposite to the applied field H. Due to the negative χ value, diamagnetic substances are said to exhibit negative magnetism. 4 Examples of diamagnetic substances are systems constituted of closed-shell atoms such as He, Ne, Ar and the rest of the noble gases; some diatomic gases such as H2 and N2; and, most organic semiconductors. In bulk form, the coinage metals Cu, Ag, and Au under study in this thesis are all diamagnetic with respective χ values of –9.7 × 10−6, –25.3 × 10−6, and –34.4 × 10−6

(unitless in SI), respectively.11

5 Figure 1.2: Scheme illustrating the arrangements of magnetic dipole moments for materials exhibiting (a) diamagnetism and (b) paramagnetism in the absence and presence of an external magnetic field H, where the length of the arrows signify the strength of the magnetic dipole. (a) A diamagnetic material is made of atoms lacking magnetic moments (spheres with no arrows); however, in the presence of an exter- nal magnetic field small magnetic moments are induced antiparallel to H (spheres with aligned small arrows). (b) A paramagnetic material is made of atoms exhibiting magnetic dipole moments in a disordered arrangement (spheres with disordered ar- rows) which are aligned parallel in the presence of H (spheres with aligned arrows). Adapted from Ref. 10.

6 1.2.2. Paramagnetism

Paramagnetic materials exhibit weak magnetism (i.e.,apositiveχ)3 due to the pres- ence of atoms (or molecules) with permanent magnetic moments; however, the inter- actions among these magnetic moments are nonexistent. 1 These magnetic moments are the result of electrons having spin or orbital moments not canceling each other out,4 a crucial precondition is the existence of permanent magnetic dipoles. 8 The lack of interactions among these localized moments and thermally-activated fluctua- tion results in disorganized magnetic moment orientations within the material, and the overall magnetization of these materials is zero in the absence of an external mag- netic field.1,4 When an external magnetic field is applied, however, there is a tendency for these moments to orient with the field H producing a net overall moment (see

Figure 1.2b).4 In many cases, thermal fluctuations at room temperature, for example, hinder total alignment resulting in a smaller overall moment due to partial alignment of the available magnetic moments.4,8 Effectively, a lower temperature will allow for a larger moment by minimizing thermal fluctuations. The relationship of the applied

field H and thermal randomization of the magnetic moments leads to the temperature dependence described by the Curie law:

C χ = (1.6) T where C is the Curie constant which is specific to a particular material, and T is the absolute temperature.12 For example, the C values for Fe, Co, and Ni are 1.23, 1.14 and 0.379 (unitless in SI).13

Based on the electrons responsible for the phenomenon, the nature of param- agnetic behaviour can be divided in two types: localized and itinerant.8 Localized paramagnetism is the result of magnetic moments by electrons found in inner or core shells of partially filled orbitals, for example, the 4f lanthanides and the 5f actinide

7 series.8 Localized paramagnetism (also known as spin paramagnetism3 or Langevin paramagnetism) is temperature dependent. 8

Itinerant paramagnetism is the result of magnetic moments by delocalized elec- trons in the valence band.8 Itinerant paramagnetism (also known as orbital paramag- netism3 or Pauli paramagnetism) is temperature independent.8,14 This temperature independence is explained by an imbalance of spin up ↑ and spin down ↓ electrons resulting from the splitting of the valence band occurring only in the presence of an external magnetic field H 2,14 as shown in Figure 1.3. In the absence of an external magnetic field H, both populations are equal.3 In the presence of H, however, there is a split of the valence band electrons2,14 due to the extra energy gained during the interactions of the electron’s spin with the magnetic field.12 For those electrons with their spin parallel to the field, the extra magnetic energy is negative and this population of electrons has a lower energy than they had in the absence of a field.12

Similarly, those electrons in antiparallel alignment acquire higher energy. Only a frac- tion of the electrons at the Fermi level opposing the field (i.e., antiparallel alignment) will be able to switch their spin and align with the field by flowing into the lower- energy freshly unoccupied orbitals found in the parallel electron population band in an attempt to equalize the energy at the Fermi levels of both electron populations, and in the process creating an imbalance.2,3,14 Another description for spin-up ↑ and spin-down ↓ electron populations is majority states describing the spin state with the larger electron population, and minority states describing spin state with the smaller population.2

Oftentimes, temperature-independent paramagnetism is similar to diamagnetism as both phenomena are the result of induced magnetic moments;2 however, param- agnetism will always result in a positive susceptibility χ, while diamagnetism will always result in a negative susceptibility value. Phenomenological, Pauli paramag-

8 Figure 1.3: Density of states for paramagnetic materials (a) in the absence and (b) presence of an external magnetic field H.Ina all electronic states are filled to the Fermi energy E F. The two spin populations N↑ and N↓ are equal in magnitude resulting in no magnetization M. However, when H is applied, there is a split of the valence band resulting in redistribution of electron populations and in the observed transient M.

netism is the result of splitting of the valence band resulting in two unequal electron

populations, while diamagnetism originates from the core electrons.

Based on our previous discussion, atoms can only exist in two forms: diamagnetic

(if comprised of completely full shells) and paramagnetic (if comprised of partially full

shells). It is only when atoms form larger materials that other forms of magnetism

are observed. This is the result of possible magnetic interactions of atoms within the

material.

1.2.3. Collective Magnetism

Thus far, the discussion has been of materials which do not present permanent mag-

netism at the macroscale as their observed magnetic behaviour is induced by an

external magnetic field H resulting in a negative susceptibility χ (diamagnetism) or

a positive susceptibility χ (paramagnetism), with both magnetic phenomena disap-

pearing once the external field is removed. There are materials, however, that show

permanent net magnetism even in the absence of an external field as a result of indi-

9 vidual atoms (i.e., atomic magnetic dipole moments) interacting cooperatively. These magnetic interactions, collectively called exchange interactions, occur among perma- nent magnetic dipoles and can lead to magnetic ordering.8 The three most common ordering schemes are ferromagnetism, antiferromagnetism, and ferrimagnetism (see

Figure 1.4). A brief description of these magnetic ordering schemes is presented in the next section.

A common characteristic of collective magnetic materials is a critical tempera- ture T below which spontaneous net magnetization is observed, and above which the collective magnetic behaviour disappears.8 Thus, the effect of temperature in mag- netic structures is decoupling of magnetic moments changing the overall magnetic behaviour from cooperative to independent behaviour.12

1.2.3.1. Ferromagnetism

In ferromagnetic substances, the permanent local magnetic moments exhibit long- range ordering arranged in a parallel fashion reinforcing one another 1,12 and the over- all macroscopic behaviour is that of a, commonly-referred, magnet (see Figure 1.4a).

It is important to note that the parallel arrangement of magnetic moments is local- ized in areas known as domains, which are present systematically through out the entire macrostructure. Magnetic domains are microscopic regions containing 1017 to

1021atoms, and within which all magnetic moments are aligned.1 The macrostruc- ture, however, shows an overall magnetic ordering resulting from vectorial parallel ordering.

Collective magnetism in ferromagnetic substances is disrupted at the critical tem- perature called the Curie point, T C. Below this transition temperature, the average values of magnetic moments are oriented such that a strong spontaneous magnetiza-

9 tion M is produced. Above the Curie temperature T C, magnetic moments decouple resulting in paramagnetic behaviour.12

10 Figure 1.4: Scheme illustrating the arrangements of magnetic dipole moments for materials exhibiting (a) ferromagnetism, (b) antiferromagnetism, and (c) ferrimag- netism in the absence of an external magnetic field H. (a) A ferromagnetic material composed of atoms with magnetic moments arranged parallel to each other. (b) An antiferromagnetic material composed of atoms with magnetic moments of the same magnitude (illustrated by arrows of equal length) arranged antiparallel to each other. (c) A ferrimagnetic material composed of atoms with magnetic moments of different magnitude arranged antiparallel to each other. Adapted from Ref. 10.

11 The typical behaviour of a ferromagnetic material in the presence of an external

magnetic field is shown in Figure 1.5a. The closed-loop trace observed in Figure 1.5a

is the result of measuring the magnetic response (magnetization M )ofthesampleas

the applied external magnetic field H is cycled. The particular MvsHloop associated with ferromagnetism is the non-linear response of M to an imposed H, and is called a hysteresis loop.15

There are several quantitative properties that can be obtained from a ferromag-

netic hysteresis loop. For instance, upon the initial application of a magnetic field

H, the magnetization M of a ferromagnetic substance will sharply increase reaching

a constant value appearing only at a sufficiently strong H.9 This maximum value is

the saturation magnetization, M s. After reaching this point, the direction of mag- netic field H is reversed during the cycle. In the absence of H, the magnetization reaches magnetic remanence, M r. Additionally, in this process at the instance when

9 M reaches zero, the applied field is called the coercive field, H c. For comparison, the magnetic responses of diamagnetic and paramagnetic substances are presented

in Figure 1.5b. The diamagnetic response is characterized by a negative slope or

negative magnetization as a result of its negative susceptibility χ. The paramagnetic

response is characterized by a positive magnetization due to its positive χ; however,

it does not present remanent magnetization M r nor coercive field H c and the signal goes through the origin (i.e.,atH =0,M =0).

Based on the characteristics of the ferromagnetic hysteretic loops, substances can

be classified as hard or soft magnetic materials. Hard magnetic materials have high H c values (10−2 to 1 T)9 resulting in a broad MvsH loop, while soft magnetic materials

−3 −2 9 15 have low H c values (10 to 10 T) resulting in narrow loops. In practice, for example, hard ferromagnets are used in permanent magnet devices such as motors

(e.g., alnicos or alloys of Al, Ni, Co, and Cu), and soft ferromagnets in magnetic

12 Figure 1.5: Simulated Mvs.H loop associated with (a) ferromagnetic, (b) param- agnetic, antiferromagnetic, and diamagnetic materials. (a) Ferromagnetic hysteresis loops contain quantitative characteristics such as saturation magnetization M s,re- manent magnetization M r, and coercive field, H c. A ferrimagnetic loop is similar to a ferromagnetic loop. For comparison, typical responses of (b) paramagnetic (at room and low temperatures), antiferromagnetic, and diamagnetic materials are also presented. See text for description.

13 recordings (e.g., supermalloy composed of 75% Ni, 20% Fe, and 5% Mo).9

As seen in Figure 1.5a, once a magnetic field H is applied on a ferromagnetic

substance, its magnetization M value depends on the intensity of H ; however, M

is also affected by the nature of the substance, by temperature, and by the history

of the sample.16 In particular, magnetization M strongly depends on the order in

which these factors are applied to the sample.16 Moreover, a time-dependence on

magnetization M is also observed. For instance, material aging can be associated

with irreversible modifications to the magnetic properties of the sample.16

1.2.3.2. Antiferromagnetism

In antiferromagnetic substances the magnetic moments are ordered in an antiparallel

fashion so that they completely cancel one another out (see Figure 1.4b). 12 The

critical temperature for antiferromagnetic materials is the N´eel point, T N.Above this temperature, antiferromagnetic materials behave paramagnetically.15

In the simplest case, an antiferromagnetic material is composed of two ferromag-

netic sublattices exhibiting the same magnetization8 intensity with opposite orienta-

tion resulting in no net magnetization in the absence of an external magnetic field

H.15 Some antiferromagnets, however, contain more than two magnetic sublattices

9 such as UO2 which contains four sublattices. Other examples of antiferromagnets are the element Cr and the metal oxide NiO. 4

For visualization purposes, an exaggerated MvsH measurement for antiferro-

magnets at room temperature is presented in Figure 1.5b. Although the trend shows

a positive response (i.e., positive slope), antiferromagnets lack macroscopic magneti-

zation and are general insensitive to the external magnetic field H.17

1.2.3.3. Ferrimagnetism

Ferrimagnetism is considered a special case of antiferromagnetism. In ferrimagnetism,

as in antiferromagnetism, at least two magnetic sublattices are aligned antiparallel

14 to each other;8 however, the strength of both moments in each population is unequal

leading to partial cancellation on the overall behaviour,12 resulting in a non-zero to-

tal magnetization (see Figure 1.4c). The magnetization M of each sublattice will be

maintained for all temperatures below the critical temperature known as the ferri-

18 magnetic Curie temperature, T C. Above the T C, ferrimagnetic materials behave paramagnetically.4 Additionally, the two sublattices may have different temperature dependencies which could lead to a particular temperature in which both magnetic moments, having an equal but opposite direction, will cancel each other – this is known as compensation temperature.19 Above or below the compensation tempera-

ture, the material will be magnetized in the direction of the applied magnetic field

M and will be non-zero,19 overall sharing characteristics of ferromagnetic materials

(e.g., similarly shape hysteresis loops).

An example of ferrimagnetism are metal oxides referred to as ferrites, with the

most famous being magnetite Fe3O4. Magnetite contains dual-valency iron atoms in a2:1Fe3+:Fe2+ mixture.15

1.3. Exchange Interactions

In previous sections, the existence of exchange interactions among atomic moments

leading to different magnetic phenomena has been been mentioned. In this section,

these interactions are discussed in more detail.

1.3.1. Exchange

In general, attractive electrostatic interactions between the negative charges of the

electrons and the positive charges of the nuclei are responsible for the cohesion of

solid materials and molecules.14 For example, for a pair of unbound H atoms at some

distance apart, electrons and protons will experience electrostatic attractive forces,

15 while repulsive forces will be present between the two electrons and between the two

protons.4 When two atoms come closer for binding, there is a tendency for redistri- bution of electrons according to their spin orientation resulting in a spin-dependent

14 coupling energy known as exchange interaction. Thus, in the H2 case, antiparal- lel spins between interatomic electrons will experience attractive forces resulting in

covalent bonding, in turn, reducing the total energy of the newly formed molecule. 4

Moreover, exchange interactions are distance dependent as they decrease rapidly with

distance.4

The exchange interactions (also called exchange forces) can lead to covalent bond-

ing which occurs when the spins of two electrons are antiparallel.14 These interac-

tions are a consequence of the Pauli exclusion principle applied to the entirety of

4 the chemical system, in this simple case the H2 molecule. In other words, the ex- change interaction between two neighbouring electrons’ spin moments on different

atoms (i.e., interatomic interactions) has the same origin as the exchange interaction

between two electrons on the same atom (i.e., intra-atomic interactions).5 However,

intra-atomic exchange interactions are larger (i.e., 90% stronger)20 than interatomic

exchange interactions.21 Pauli’s exclusion principle states that two electrons can have

the same energy only if they have opposite spins4 (i.e., two electrons in the same

system cannot have identical quantum numbers).14 Finally, exchange forces mainly

depend on interatomic distances and not on crystallinity as shown by the existence

of amorphous ferromagnets.4

The term exchange in exchange interactions is derived from the fact that elec- trons are indistinguishable (due to Heisenberg’s uncertainty principle) and electron

4 exchange takes place between atoms. In our H2 example, electron a moves about proton b and electron b moves about proton a exchanging at a rate of 1018 times per

second.4 The process of electron exchange requires exchange energy. The exchange

16 energy interacting between atoms a and b, for example, can be described by the

Heisenberg model as follows:4,5,14,22

Eex = −2Ja·bSaSb (1.7)

where E ex is the exchange energy of interaction between atoms, J a·b is the particular

exchange constant for atoms a and b,andS a and S b are the spin angular momentum

for electrons in atoms a and b respectively. The exchange constant J (also called

exchange integral)2 is related to the overlap of the charge distributions of the atoms

considered,14 and it is a function of the interatomic distance between the atoms.5,22

The interatomic distance dependence of the exchange interaction is responsible for

several effects discussed in the next subsections.

The exchange constant J is the key to understanding the magnetic phenomena

described here. For instance, J can acquire values above or below zero. When

J > 0 the electrons’ spin pair ferromagnetically or parallel; conversely, when J < 0

the electrons’ spin couple antiferromagnetically or antiparallel 22 as can be seen in

Figure 1.6. In order to understand the parallel pairing of electrons while maintaining

the Pauli exclusion principle intact, the symmetrization postulate is required.

Electrons, as fermions, are identical particles which, in an indistinguishable situ-

ation, can only be found in antisymmetric states23 described by the symmetrization

postulate. The symmetrization postulate of an electron contains a spatial wavefunc-

tion and a spin wavefunction2 with one of them being in an antisymmetric state

when electrons are coupled. In antiferromagnetism, the space functions are symmet-

ric allowing for the electrons to be close resulting in antiparallel spin alignment, for

overlapping orbitals, and for a strong electron-electron repulsion.2 In ferromagnetism,

the space functions are antisymmetric resulting in distant electrons with less orbital

overlapping which allows for parallel spin coupling.2

17 Figure 1.6: The Bathe-Slater plot illustrating the strength of direct exchange coupling and its effect on the exchange constant J as a function of interatomic distance (r a·b) and the radius of the incomplete filled d-shell orbitals (r d) of an atom in a material. The interactions lead to ferromagnetism (↑↑)orantiferromagnetism(↑↓) magnetic ordering. Adapted from Ref. 24.

1.3.2. Direct Exchange Interaction

The direct exchange interaction arises from non-negligible overlap between the mag- netic orbitals of two adjacent atoms, and occurs between two neighbouring spin mo- ments.5,22 Since there is orbital overlapping, a correlation exists between the nature of the exchange interaction and the interatomic distance of the atoms considered. This correlation, discovered by Slater, states that the nature of the exchange interaction

(whether positive or negative interaction) corresponds to the ratio r a·b / r d,where

r a·b represents the interatomic distance and r d the radius of the unfilled d-shell or- bitals.5 A graphical representation of this effect is the Bethe-Slater curve presented

in Figure 1.6.

It is important to note that direct exchange as a basis for ferromagnetism is

only applicable to a few cases, for example CrO2 and CrBr3, due to their short-

18 ranged interactions.22 The 3d transition metals Fe, Co and Ni all exhibit long-range interactions. See Section 1.4.1 dealing with conventional magnetism.

1.3.3. Indirect Exchange Interactions

Unlike direct exchange, in which there is a direct overlap of electrons wavefunctions of adjacent atoms without an intermediary atom, in indirect exchange the wavefunction of a magnetic atom overlaps with the wavefunction of a nonmagnetic intermediary atom or ligands which in turn overlaps with the wavefunction of a second magnetic atom.2 There are several forms of indirect exchange interactions: superexchange, double exchange, and Ruderman-Kittel-Kasuya-Yosida commonly refer to as RKKY exchange.2 These interactions are described next.

1.3.3.1. Superexchange

The name superexchange is derived from its relative longer range interaction com- pared to the short-ranged direct exchange.2 Consider for example the oxide NiO,

which is an antiferromagnet. As a crystalline solid, Ni–O–Ni atomic arrangements

(i.e., 180◦ angles) are found through out the structure allowing for the 3d Ni orbitals

to overlap with the 2p oxygen orbitals,12 as can be seen in Figure 1.7. Thus, in our

arrangement, the magnetic Ni atom positioned at the left will have its ↑ electron

paired with the ↓ electron of the 2pz orbital from the nonmagnetic O atom. The ↑

2pz electron in the O atom will pair with the ↓ electron of the Ni atom positioned to the right of the O atom. The net result is that adjacent Ni atoms have opposite spins, rendering an antiferromagnetic structure.12 Therefore, superexchange interac-

tions lead to antiferromagnetic coupling of the metal atoms by orbital overlapping

with the O atom, which remains nonmagnetic.2 Additionally, the same interactions

can be found in ferrimagnetic materials.5

19 Ni O Ni

3dz2 2pz 3dz2

Figure 1.7: The superexchange interaction found in NiO by the overlap between Ni 2 d z and O pz orbitals. Adapted from Ref. 12.

Figure 1.8: The double exchange interaction found in Fe3O4. An O atom (red sphere) connecting two octahedral Fe ions allows for electron hopping (blue sphere) between Fe2+ and Fe3+ and vice versa. The dashed arrows labeled 1 and 2 are indicating the two-stage nature of the overall process. Adapted from Ref. 2.

1.3.3.2. Double Exchange

The process of double exchange occurs in ions exhibiting mixed valence configurations leading to ferromagnetic arrangement.8,15 This interaction arises between 3d ions

having electrons hopping from an ion core to the next retaining their spin direction

in the process.15 Unlike superexchange, which leads to antiferromagnetic ordering,

in double exchange interactions, the indirect coupling of spins accross a diamagnetic

atom leads to ferromagnetic ordering.2 An example of double exchange is the oldest

known magnetic material Fe3O4 magnetite, which contains Fe atoms in two different valence states Fe2+ and Fe3+ that are ferromagnetically coupled.2 The process of

double exchange in magnetite is presented in Figure 1.8.

20 The name double exchange derives from the fact that two metal atoms (differing

in valency by one electron) in the system are bonded through an O atom creating a

bridge by which this “hopping” electron moves back and forth between the two metal

atoms.2 It is accepted that the process of double exchange occurs in two stages as seen in Figure 1.8. In the first stage an incoming electron from the lowest oxidation metal species moves to the bridging O atom; however, the O orbitals are full resulting in the expulsion of a new electron. Both electrons retain the same spin as spin flip is forbidden. In the second stage, the new electron hops from the O to the other metal species. The entire mechanism is a rapid process resulting in a partial delocalized negative charge.

1.3.3.3. RKKY

The acronym RKKY stands for Ruderman and Kittel (1954), Kasuya (1956), and

Yosida (1957) whom have been acknowledge as the discoverers of this particular phe- nomenon.15,20,21 RKKY interactions are described by the fact that a single localized

magnetic impurity can create a nonuniform-oscillating spin polarization mediated by

conduction electrons of the host material inducing areas of antiparallel and parallel

configurations depending on the separation distance near the magnetic impurity.15

The RKKY effect is illustrated in Figure 1.9. An analogy to this phenomena are the

waves created by a stone thrown into water creating ripples that travel through the

medium.20

RKKY interactions are thought to play a role in the magnetism observed in lan-

thanides or rare-earth metals (e.g., Gd). In the case of lanthanides, the localized

moments in their 4f core orbitals are believed to be capable of spin polarizing the

conduction electrons of the outer 6s shell.14,21

21 Figure 1.9: The RKKY effect illustrated by the exchange parameter J as a function of the interatomic distance r a·b. The interactions lead to ferromagnetism (↑↑)or antiferromagnetism (↑↓) magnetic ordering. Adapted from Refs. 24 and 25.

1.4. Conventional and Unconventional Magnetism

Thus far the discussion in previous sections has been of magnetic moments present in well-defined atomic sites, like those founds in transition-metal oxides and rare-earth metals.21 The iron-series transition metals, however, do show magnetism caused by

delocalized electrons.21

1.4.1. Conventional magnetism

In this thesis, conventional magnetism refers to the phenomenon observed in the 3d

transition metals Fe, Co, and Ni all of which are ferromagnets at room temperature.26

For completion, the fourth and last ferromagnetic element in the periodic table at

just below room temperature of 16 ◦C4 is Gd belonging to the lanthanides or rare-

earth metals, while the rest of the elements in that series behave paramagnetically as

pure metals at room temperature.2 In the 3d transition metals, both the 4s and 3d

22 electron bands contribute to the density of states at the Fermi level.27 However, only

the 3d electrons contribute to the ferromagnetic characteristics while the 4s electrons

dominate the conduction characteristic due to their higher mobility compared to their

3d counterparts.27 Essentially, the electrons closer to the nucleus (i.e., lower shells

such as n=1, 2, and some 3 orbitals) are bound by the nuclear charge and form an electrostatic potential in which the outer electrons move, in this case the delocalized electrons originate from the 4s and 3d orbitals.2,20

The type of ferromagnetism found in the transition metals Fe, Co and Ni is known

as band or itinerant magnetism.28 Their magnetic behaviour derives from an im-

balance of spin-up ↑ and spin-down ↓ 3d electron populations due to strong direct exchange between them,26,27 as depicted in Figure 1.10. This imbalance leads to un-

paired electrons and a net magnetic moment.26 The rest of the 3d transition metals

feature identical 3d electron populations of both spins (i.e., equal number of elec-

trons with spin-up ↑ as with spin-down ↓) producing no net magnetization.27 The

alignment of spin momenta is a competition between the exchange energy and ki-

netic energy effects.24 If the alignment is parallel (i.e., ferromagnet), there is a gain

in exchange energy and a loss of kinetic energy; while the opposite is true for non-

magnetic materials.24 The Stoner model explains the reason why some 3d transition

metals are ferromagnetic and others are not: the product of the density of states and

the exchange integral must be greater than unity for spontaneous spin ordering to

emerge – this is known as the Stoner criterion.29 Mathematically, the Stoner criterion

is defined as follows:

D(F) · J>1 (1.8)

where D(F) is the density of states at the Fermi level and J is the exchange integral for ferromagnetism.30

23 Figure 1.10: The Stoner model for ferromagnetic transition metals. Only the density of states of the 3d shell is shown. Filled states (electrons) below the Fermi energy E F are shaded, while empty states (holes) are shown blank. Symbols ↑ and ↓ represent the spin of the electron populations. Adapted from Ref. 2.

Before continuing with this discussion, it is necessary to define magnetic moment units. The units for magnetic moment, μ,istheelectromagnetic unit of magnetic moment or, unofficially, emu which can be further defined as erg/Oersted.4 In most

cases, the magnetometry data in this manuscript will be presented as emu per unit

mass. However, as described earlier by Equation 1.2, when dealing with small volumes

such as atoms or unit cells, atomic magnetic moments are given in units of Bohr

−21 magnetons, μB, where 1 Bohr magneton is equal to 9.27 × 10 erg/Oersted in CGS units4 or 9.27 × 10−24 J/T in SI units.1

The Stoner model of delocalized electrons in the iron transition metal series is

supported by the fact that the experimental average magnetic moment per unit vol-

ume is less than that expected of the constituent atoms; moreover, these values are

fractional rather than whole numbers.9 For instance, the experimental value (from

2 saturation magnetization M s) for metallic iron is 2.216 μB/atom. This is the re- sult of electron delocalization reducing the spin magnetic moment to 30–50% of their

atomic values,31 in addition to the crystal field effects quenching the orbital magnetic

24 moment to a total of 5%31 and 10%.2

The condition for the orbital moment quenching is the result of the electric field

(i.e., crystal field) produced by the surrounding atoms in the solid.4 This field has the

symmetry of the crystal involved4 and the orbitals are locked into a specific direction

by the bonding anisotropy of the lattice.2 Moreover, the orbit is strongly coupled to

the crystal lattice which prevents the orbital moment from turning (i.e., anisotropic)

in the direction of the magnetic field when the specimen is exposed to an external

magnetic field H.2,4 On the contrary, the spin moment is able to align (i.e., isotropic)

with the external magnetic field H due to its weak coupling to the lattice and to the

orbital moment.2,4 The net result is that mostly the spins contribute to the overall magnetic moment of the material.4

1.4.2. Unconventional magnetism

Unconventional or unexpected magnetism appears in some diamagnetic materials

when their dimensions are reduced due to the emergence of quantum effects at the

nanoscale.24 In general, surface atoms of materials have fewer neighbours compared to

bulk atoms resulting in the former having a lower coordination number and unsatis-

fied bonds (i.e., are less stable than bulk atoms).32 Physically, reduced dimensionality

results in nanostructures having a larger fraction of their atoms at the surface 32 and

to the appearance of new surface and interface states as a result of higher surface-to-

volume ratios.24 Additionally, the appearance of point defects such as vacancies or

interstitial atoms, could potentially play a role in inducing magnetism in the nanoma-

terials.24 Vacancies, for example, induce the formation of localized electronic states

creating finite magnetic moments which can mediate magnetic interactions. 24,33

An example of the strong connection between magnetic behaviour and small di-

mension is the experimental work with Cr and Rh clusters. In the bulk Cr is the only

element to show antiferromagnetism at room temperature.34 However, small clusters

25 of 20-133 atoms show a large magnetic moment shifting from 0.62 μB/atom in the

34 bulk to almost double at 1.16 μB for Cr58. Similarly, for the paramagnet Rh, 9-atom

clusters show a magnetic moment of 0.8 ± 0.2 μB/atom, Rh34 a value of 0.16 ± 0.13

μB,andatRh60 and larger size clusters the moment reverts back to zero which is its bulk characteristic35 in the absence of an external magnetic field H.

1.5. Unconventional Magnetism in Coinage Metals

While nanodimensionality is believed to be a factor in the observed magnetism in

coinage and noble metals, there is evidence that other factors affect such behaviour.

At the nanoscale, recent studies suggest that nanodimensionality induces magnetism

in intrinsically diamagnetic bulk materials.36,37 It is generally accepted that the mag-

netism observed in NPs of coinage metals is attributed to the passivation of their

surface. For instance, the first report of magnetic behaviour in elemental Au was pub-

lished in 1999 and included 2.5 nm Au NPs coated with poly-N-vinyl-2-pyrrolidone

resulting in a saturation magnetization M s of ∼1.8 emu/g at 4.2 K and 5 T (with an extrapolated magnetic moment of ∼1.3 emu/g at 1 T) by SQUID magnetometry.37

Since then, most research in the area has focused on the anomalous magnetic be-

haviour of Au NPs and films coated with strong binding alkanethiols – albeit without

a clear understanding of the phenomenon.26 Several magnetic behaviours have been

reported in the literature for different Au NP sizes capped with dodecanethiol. In

1.4 nm particles, the behaviour has been described as ferromagnetic, measuring a

38 saturation magnetization M s of ∼0.4 emu/g; while for particles of 5 nm (M s of ∼5 × 10−3 emu/g, taking note of the orders of magnitude) and 12 nm in size their

behaviours have been reported as superparamagnetic and diamagnetic, respectively39

– all recorded at 300 K and 1 T by SQUID magnetometry. Based on these results, it

is clear that the observed magnetic behaviour is dependent on the NP size; however,

26 it is not clear why this is the case.

Initially, most of the reported magnetic studies regarding the coinage metals were

of Au NPs covered with a SAM of a capping agent. However, later studies expanded

to dodecanethiol-covered Ag and Cu NPs.26,38,40 For example, the same study found

dodecanethiol-covered NPs of 1.9 ± 0.2 nm Au, 2.3 ± 0.3Agnmand2.3± 0.2

nm Cu to have saturation magnetization M s values of ∼4.8, ∼5.4 and ∼1.3 emu/g, respectively, from SQUID magnetometry at 300 K and 6 T – their extrapolated mag- netization values at 1 T were almost identical as those at 6 T. 40 All three metals

are isoelectronic having a d10s1 valence electronic configuration and are well-known

to be diamagnetic as a result of delocalization of the free electron density in the s

band.2,30 In general, it is believed that the nature of the capping agent plays a role in

the magnetic behaviour ultimately observed in these metals.26 For instance, Au NPs

capped with a weakly interacting agent such as tetraalkyl ammonium bromide results

in diamagnetic characteristics.26,38 While a SAM of a strongly interacting agent, such

as thiols, results in ferromagnetic characteristic.26,38 This ferromagnetic behaviour is

believed to be the result of charge transfer between the Au and S atoms generating

unoccupied orbitals in the 5d states of Au.26,36,38 In essence, the dipole inherent in

the capping thiol along with the electron charge transfer induces a surface dipole at

the interface of the nanostructure.26 Experimentally, it is observed that the surface

plasmon resonance (SPR) signal (which corresponds to the collective oscillation of

electrons within the particle) at 550 nm for 1.4 nm diameter Au NPs is absent when

compared to the same NPs capped with a weakly-bound surfactant.38 This observa-

tion suggests that the Au-S bond induces the 5d Au electrons to become localized or

partially localized at the surface.38 It is then argued that the presence of a SAM at

the surface of a NP can modify the electronic structure of the metal. 26,41

A model proposing to explain the unusual magnetic behaviour in SAM-covered

27 noble metals is the one put forth by Hernando et al.42,43 This model utilizes the lit- erature’s experimental observations obtained with SAMs formed on both NPs and thin films.42,43 When SAM domains are formed on thin metal films, they become a potential well (due to highly ordered and densely packed capping agents) where sur- face electrons are captured; thus, these captured electrons are confined to movement at the surface plane. Additionally, the energy of these electrons is minimized when their interactions align their spins in a parallel arrangement. (The idea is analogous to Hund’s rule in which the electrons promote orbital degeneracy by not pairing until all sub-orbitals are half filled.) This model is able to explain several experimental observations. For instance, large areas covered by SAMs will create large magnetic moments as observed in thin films compared to small magnetic moments in NPs.

From this model, it would be expected that larger NPs would generate larger mo- ments. This is not the case however, as the contributions of the core electrons will eclipse those of surface electrons. The absence of the SPR effect in NPs is also ex- plained by this model as electron localization at the surface does not permit oscillation as expected for SPR.42,43 Conversely, scientific literature of Au NPs exhibiting both

SPR and ferromagnetism has been reported in 6.7 nm Au NPs coated with both oleic acid or oleylamine capping ligands showing an optical absorption at 528 nm and a

−2 saturation magnetization M s of ∼1.45 × 10 emu/g and a coercive field H c of 5 mT at 300 K and 1 T by SQUID magnetometry.41

Moreover, recent evidence suggests that magnetism in coinage metal nanomate- rials can be obtained without capping ligands adorning their surface. For instance,

2.5 nm Au NPs prepared in solution with tetrakis(hydroxymethyl)phosphonium chlo- ride (THPC) as a reducing agent showed ferromagnetic behaviour with THPC at the surface, and after several rinses with HCl in order to remove the capping agent.36

The values for saturation magnetizations M s were ∼0.1 emu/g for THPC-capped,

28 and half of that value at ∼0.05 emu/g for HCl-washed particles. An issue with so- lution synthesis as a preparation approach is that although no evidence of THPC was observed in spectroscopy studies after washing the NPs HCl, trace amounts of undetectable THPC could still be present at the surface.36 To eliminate the possi-

bility of residual capping ligand inducing an effect on the ferromagnetic response,

the same research group prepared fresh Au NPs photochemically in solution. The

resulting Au NP population in the range of 5-30 nm in diameter in the absence of

capping agents (with only acetone and isopropyl alcohol present) resulted in a satu-

36 ration magnetization M s of ∼0.04 emu/g at 300 K and 1 T. It is worth noting the large polydispersed population on the photochemically synthesized NPs is the result

of lack of size management in the absence of a capping ligand.

More evidence of surface-unmodified Au behaving ferromagnetically was also ob-

served in cluster-deposited films.44 In this case, Au clusters of ∼2.6 ± 0.6 nm in size

(measured by cluster beam deposition, although scanning electron microscopy mea- sured them as ∼10-30 nm) used to make films of 28 and 175 nm in thickness resulted

3 in saturation magnetizations M s of ∼10 and ∼1emu/cm and coercivity ranges H c of 5.7–10 and 7.3–11.1 mT, respectively, at 300 K and 1 T by SQUID magnetome-

try.44 The difference in magnetic characteristics observed in both films was attributed

to more material present in the thicker film which would result in a more bulk-like

behaviour.44 Additionally, since both films were created using the same cluster diam-

eter, higher fluence deposition reduces the distance among the clusters which favours

coalescence resulting in bulk-like magnetic behaviour.44

In addition to capped- and, paradoxically, uncapped-nanomaterials behaving fer-

romagnetically, a third process that induces magnetism to coinage nanoscale struc-

tures is oxidation. It has been shown that solution-synthesized ∼26-nm CuO NPs

◦ air-annealed at 800 C exhibit room-temperature ferromagnetism (M s of 0.018 emu/g

29 45 and H c of 5.7 mT at 300 K and 1 T). The magnetic behaviour was attributed to the presence of oxygen vacancies at the surface.45 The presence of oxygen deficient sites re-

sult in 3d localized electrons in the metal ion which changes its oxidation state from

Mn+ to M(n−1)+,46 which could potentially allow for double exchange interactions.

The same research team demonstrated ferromagnetism in bulk CuO/Cu2Ocompos-

47 48 ites and CuO/Cu2O microspheres attributed to an indirect double exchange in- teraction at the interface between the two oxides.47,48 The saturation magnetizations

M s for the bulk composite and microsphere systems were 0.04 emu/g (annealing at 900 ◦C for 2 hr) and 0.00125 emu/g (treatment at 180 ◦C for 12 hr), respectively, at

300 K and 0.8 T by SQUID magnetometry.47,48

Apart from oxide interfaces, non-stoichiometric oxidation is another approach to

achieve ferromagnetism. Theoretical studies of defect formation in a 48-atom Cu2O supercell indicates that an extra interstitial O atom surrounded by six or four Cu

atoms induces ferromagnetic behaviour.49 Despite the supercell consisting of a 2:1

Cu:O stoichiometric ratio, the proposed defects result in a 6:1 and 4:1 Cu:O atomic ra-

tios respectively.49 Other computational studies predict ferromagnetism at 4:5, 16:15

and 28:27 Cu:O atomic ratios.50

In the Ag case, computational studies show that Ag surfaces exposed to O2 can result in metastable ferromagnetic interactions.51 Specifically, Ag(111) surfaces are

able to dissociate O2 molecule as they approach the surface in a vertical trajectory interacting with the top of a Ag atom or with an fcc-hollow site. A surface cover-

age consisting of 0.5 monolayer O onto Ag (i.e., a 2:1 Ag:O ratio considering only

the first Ag atomic layer) creates magnetic states due to a combination of direct ex-

change interactions between the 4d orbitals of Ag and its nearest O 2p states; and

superexchange interactions between neighbouring O atoms via the Ag atom.51 These

particular direct exchange interactions result in ferromagnetic behaviour, while the

30 superexchange interactions result in antiferromagnetic behaviour, with the former being stronger than the latter.51 Moreover, a second form of indirect exchange in- teractions in the form of RKKY are also present among the O atoms at the surface.

These RKKY interactions are affected by the distance among O atoms (i.e., depend- ing on the O coverage) with Ag:O ratios of 9:2, 4:1 and 3:1 resulting in non-significant

RKKY interactions.51 Finally, increasing oxidation near the metal surface by intro- ducing sub-surface O to their theoretical studies induced these O atoms to withdraw electrons from the Ag surface, quenching the ferromagnetic behaviour. 52

In their totality, the wide variety (and often contradictory) experimental evidence and theoretical predictions describing scenarios in which ferromagnetism is induced in coinage metals is exceptionally complex, which may indicate that there are yet unresolved principles at work that warrant further study.

1.6. Nanomaterials

The boom in nanoscience and nanotechnology has been achieved by technological advancements allowing for the tailored fabrication of nanomaterials which in turn provide the building blocks for both fields.53 In these materials, nanodimensionality is achieved by restricting one or more dimensions to the nanometer scale. 53 It is generally accepted that in order to be considered a nanomaterial, the scale of one or more dimensions must be ≤ 100 nm. By restricting different dimensions, nanostructure materials are classified into three categories: thin films in which only one of the three dimensions is of nanometer scale, nanowires which have two of the three dimensions of nanoscale, and nanoparticles in which all three dimensions are of nanometer scale.24

The work presented in this thesis focused on nanoparticles and thin films. As such, nanowires will not be discussed further.

31 1.6.1. Nanoparticles

Regarding the iron-series ferromagnetic elements (i.e., Fe, Co and Ni), nanodimen-

sionality also affects their magnetic behaviour. Specifically, clusters Ni150,Co450,and

Fe550 (equivalent to three layers for Ni, and to approximately four to five layers for Co and Fe) demonstrate bulk-like ferromagnetic behaviour, while smaller clusters experimentally show a higher magnetic moment per atom when compared to their bulk counterparts.54 The non-linear decaying trend for this particular ferromagnetic behaviour (i.e., from high moments in small clusters to bulk-level moments in larger structures) is attributed to the dominant electronic effects caused by surface atoms. 54

A property closely related to NP size, specifically volume, is superparamagnetism.

Superparamagnetism is characterized by a fluctuation of the magnetic moment μ of

the NP as a whole while maintaining the ferromagnetic ordering of the individual

atoms.55 Magnetometry measurements of superparamagnetic materials result in an-

hysteretic, but sigmoidal, MvsHcurves55 at room temperature,56 with similarities to

localized paramagnetic measurements but with higher magnetic moments per particle

reaching 105 times the atomic moment.57

An interesting aspect of coinage metal NPs is their unusual temperature de-

pendence to their magnetic behaviour compared to ferromagnetic metal NPs. As

previously described, the magnetism of conventional ferromagnetic NPs (and bulk

8 magnetic materials) follows the Curie temperature, T C. Above this critical tem- perature a ferromagnetic material becomes paramagnetic, if sufficiently small these

NPs will behave superparamagneticly at room temperature. For example, in 16 nm

58 CoFe2O4/SiO2 nanocomposite NPs, ferromagnetism is lost at 127 K. On the con- trary, the unexpected behaviour of thiolated 2.5 nm Au NP is that they retain their

ferromagnetic characteristic even at 300 K.36

32 1.6.2. Thin films

Conventionally, studies of magnetic thin films are performed by growing 3d magnetic transition metals, for example, onto nonmagnetic substrates.24 An inherent problem with this approach is the strain the thin films are subjected to due to the potential lattice mismatch with the substrate which could influence the magnetic properties of the film under study;24 thus, the lattice parameters of a film will differ from those of the bulk.15 In thin films the intrinsic magnetic properties (e.g., magnetization, Curie temperature, and anisotropy) may differ greatly in thin films compared to bulk scale; however, this influence is mostly restricted to the first atomic layers at the interface.15

This is believed to be the case as sufficiently thicker films adopt equilibrium lattice parameters (aided by atomic-scale dislocations) far from the substrate.15 Additionally, the crystal structure in nanocrystals is determine by factors such as lattice energy, surface energy, and dislocations.12 For example, Cu in the bulk is face-centered cubic but can be found as a body-centered cubic in thin films.12

The importance of magnetic thin films as a study subject is the technological convenience and improved performance in microelectronics.9 The first real application of nanotechnology came in the form of magnetic films in applications on read-out heads in hard drives due to the giant magnetoresistance phenomenon.59

1.7. Sputtering process

The process of sputtering was utilized to synthesize all nanomaterials presented in this thesis; thus, a description of the process is provided here. The phenomenon of sputtering was first demonstrated in the laboratory by William Robert Grove in

1852.60 Sputtering is defined as the removal of material from objects (i.e.,erosion) by kinetic energy transfer in collisions of energetic atomic projectiles.60,61

As a technique, sputtering, also known as sputter etching, is used for a number of

33 applications which require careful, microscopic erosion of a surface such as patterning, cleaning, and depth profiling.62 Additionally, it can be used as an analytical method in which an ion beam is used as the probe, for example, in secondary ion mass spectrom- etry, ion scattering spectroscopy, direct recoil spectrometry, and mass spectrometry of recoiled ions.61 Lastly, collection of the atoms produced from the blasted surface onto a substrate is referred to as sputter deposition. 62 Currently, sputter deposition, as part of the physical vapour deposition techniques such as thermal evaporation and pulsed laser ablation, is used extensively in the scientific community and industry to deposit thin films on substrates.63

1.7.1. Sputtering theory

Sputtering is a purely physical process consisting of ion bombardment onto the target material resulting in ejected source atoms as vapour (i.e., sputtering).63 In almost all cases, sputtering is achieved with ion bombardment relying on the transfer of momentum and kinetic energy from the incident particle to the surface atoms. 62 It is important to note that the ion species used for bombardment is independent of the particle’s charge;61 however, ions must contain enough kinetic energy in order to be able to break the bonds of the source material. The use of ions for sputtering is the result of easier maneuvering compared to neutral atoms. For instance, ions respond to electric fields and potential biases as a consequence of their charge, while atoms do not.62

In general, there are two systems used to generate ions in sputtering: ion beams and plasmas.62 In ion beam systems, ions are extracted from an ion source and aimed towards the target.62,63 Both, the ion source and the target, are physically separated from one another. In this case a negative voltage at the cathode (i.e., either the target or positioned behind the target) attracts the positively charged ions (usually

Ar+) with sufficient energy to initiate sputtering.63

34 In contrast, in plasma systems, the surface of the target itself is immersed in

the plasma.62 In the simplest experimental arrangement, during a plasma sputtering

deposition, a cathode and an anode are positioned near one another in a chamber

under vacuum conditions.63 Then, Ar is introduced into the vacuum chamber before a

voltage is applied between the electrodes, in the process igniting a glow discharge (i.e.,

plasma).63 The glowing plasma is due to the ionization of the gas atoms conducting electricity between the two electrodes.62 The Ar+ ions forming the plasma have a

short mean free path as a result of electrons experiencing quantum tunneling between

Ar atoms and Ar+ ions at close distances of approximately a few A˚ in the process reversing roles (i.e., the ion becomes an atom, and the atom is ionized).63 The process

is called a symmetric charge exchange collision.63

Magnetron sputtering is a popular variation of plasma systems used in 95% of

all sputtering applications.62,63 A magnetron works by using a static magnetic field

located parallel to the cathode surface (i.e., the target or source material).62 Ion

bombardment not only releases source atoms but electrons are also emitted from the

target.62 These secondary electrons are trapped by the magnetic field to move in a direction parallel to the target surface (i.e., perpendicular to the magnetic field and the electric field).62 The magnetic field forms a current loop of drifting (i.e.,

low kinetic energy) secondary electrons trapped just above the target forming the

plasma.62 Compared to other deposition techniques, magnetron sputtering possess

several advantages such as high deposition rates, ease of sputtering any metal resulting

in high-purity films, high adhesion of films onto substrates, ability to coat heat-

sensitive substrates, and excellent film uniformity on relative large areas.64

1.7.1.1. Sputtering yield

The ability to measure sputtering rate, or the rate of sample etching, which is defined

as the amount of material removed per sputtering time, is usually proportional to

35 the thickness of the deposited film.61 However, a more accurate method to describe

the sputtering phenomenon is known as the sputtering yield Y defined as the atomic

ratio of sputter ejected atoms per incident particles,61–63 and has been theoretically

modeled by various authors.63

Experimentally, the sputter yield Y, in units of source grams per incident ion, is calculated as follows:61

V ρ m e Y (g/incident ion) = crater metal a (1.9) ItNA

where V crater represents the quantity of material removed, ρmetal is the density of

sputtered material, ma is mass of the sputtered species, e is the charge of the electron (note that for noble-gas sputtering, an ion is generated per electron so that e is 1.6 ×

−19 10 C), I is the ion current, t is the sputtering time, and N A is Avogadro’s number. In general, sputter deposition rate scales with incident ion energy (applied volt-

age between the electrodes) and with an increase in ion current flux (Ar flow rate).62

Specifically to nanoparticles, changing plasma conditions (e.g., by adjusting the mag-

netic field distribution at the surface target or by adjusting gas flow) results in

nanoparticles becoming well-crystallized with faceted morphology.65

1.7.2. Gas-phase nanoparticle generation

It has been observed that nanoparticles can be formed in the gas phase while sput-

tering. Synthesis methodology, alongside material preference, is the fundamental

step that defines the characteristics, properties and, in most cases, the application

of nanoparticles. The general strategy for all synthetic procedures involving the gas

generation of nanoparticles is the rapid increase in concentration of vapour chemical

species followed by rapid cooling leading to nucleation and particle growth. Gas-phase

preparation of nanoparticles is able to separate nucleation and growth stages tempo-

36 rally and spatially.65 These two mechanisms along with aggregation are described as follows:66,67

• Nucleation: high vapour concentrations of atomic species undergo

collisions inducing cluster formation. These clusters are composed of

only a few atoms and form the nuclei which can further gain mass in

the growth phase.

• Growth: once nucleated clusters are formed, particle growth occurs

via two processes: surface growth and coalescence. In surface growth,

atomic species are added to the surface of formed clusters increasing

their size in the process. During coalescence, nucleated clusters at

high concentrations collide and combine forming larger particles. Both

mechanisms can occur simultaneously (at different rates) or separately

depending on the conditions. In both instances, particle growth ceases

when a critical size is reached and further growth becomes unfavourable.

• Aggregation/coagulation: unlike coalescence, fusion of particles does

not occur in aggregation. Aggregate complexes are formed by several

individual nanoparticles in an attempt to reduce their high surface en-

ergy without fusing into larger structures.

Gas-phase nanoparticle generation was pioneered by Gleiter et. al. in 1984.67 In their work, they were able to synthesize 6 nm Fe nanoparticles by sublimating Fe powder under He gas conditions.68 This methodology is called inert gas condensation

(IGC). Since then, IGC has been modified in several ways. For example, vapour- ization could be accomplished with lasers (i.e., laser ablation). Also, reactors have moved from single-chamber to multi-chambers and tubular reactors allowing for bet- ter particle sizing control.

37 In IGC, the first step is the application of thermal energy to a bulk material

with a tungsten heater within a vacuum chamber in an inert gas atmosphere (usually

He or Ar gas).66 High temperatures are required for vapourization and may reach

1700 ◦C.69 Due to the high temperatures utilized in this process, the source materials

are limited to metals. The source vapour moves away from the “hot” area due to

thermophoresis (i.e., temperature gradient) and into cooler areas within the system

such as cold fingers.69 The role of the inert gas is to influence the collision frequency

of the vapour species and to remove their thermal energy.67 For example, collisions

between metal atoms and inert gas atoms lead to kinetic energy transfer from the

metal to the gas. Conversely, collisions between metal atoms lead to nucleation.

An advantage of a multi-chamber plasma sputtering set-up (as the one utilized in

this research) over IGC is the control over factors affecting particle size. For example,

an increase in incident ion energy ion current flux directly affects the amount of target

atoms released in the system such that a higher atom concentration induces larger

nanoparticles as result of more collisions. Also, the inert gas flow rate affects the

residency time of nucleated clusters growth. For instance, high flow rates mobilized

the clusters faster reducing growth time resulting in smaller particles than at lower

flow rates. Third, the distance traveled by the clusters inside the system the longer

the distances, the more chances for collisions resulting in larger nanoparticles.

1.8. Research goal

Several nanomaterials exhibit unique physical properties due to their nanoscale di-

mension which overrules the critical lengths that characterize conventional physical

phenomena.70 For instance, non-magnetic bulk materials expressing magnetic be- haviour at the nanoscale is a recent phenomena that has caught the attention of many researchers.71–76 The origin of this “unconventional magnetism” is currently

38 unknown and is being debated with several, and often contradictory, mechanisms.26

Much of the scientific literature on this topic encompasses nanomaterials made of noble and coinage metals, such as Cu, Ag, and Au, which are diamagnetic (i.e., non-magnetic) in the bulk.77 Additionally, most of this research demonstrates that in order for this unconventional magnetism to appear, ligands forming a self-assembled monolayer (SAM) must be present at the surface of the material.38,40,78 However, some others have shown that SAMs are not required.44 Yet again, others have shown that magnetism ensues when a specific oxide species, whether an oxide expressing a non-stoichiometric ratio with the metal 49,50,79 or a mixture of two different oxide valencies,45,47,48 are required to either observe or maximize the magnetic signature.

The main scientific quest in nanomagnetism is to create, explore and understand new nanomagnetic materials and phenomena. 80 Thus, the main goal of this work is to study how unconventional magnetism in nanomaterials is affected by three factors: the elemental composition of the material when synthesizing structures of Cu, Ag, and Au; the size of nanostructures by synthesizing nanoparticles and thin-films; and, surface modifications by studying pristine (i.e., ligand-free) vs oxidized nanomaterials.

In order to study all three effects independently of each other, the gas-phase synthesis via sputtering method was adopted.

As previously stated, all three coinage metals studied are diamagnetic in the bulk.

All three metals belong to Group 11 of the Periodic Table and, thus, isoelectronic

(referring to the number of valence electrons). They exhibit a d 10 s1 configuration; thus, an independent unbound atom will exhibit magnetic behaviour due to its un- paired s-orbital electron. However, as atoms combine to form structures this magnetic behaviour is disrupted as electrons become delocalized resulting in the well-known conductive characteristic of metals. Our interest resides in our ability to synthe- size nanostructures of varying sizes within which bulk properties may not be fully

39 established and where size effects on their magnetic responses may be observable.

The surface (whether modified or not) of nanomaterials plays an important role

on their properties. Physically, this is the result of an unusual quantity of surface

atoms vs volume atoms inside the bulk of the material. For example, a metal cluster consisting of 13 atoms (or first complete shell) contains 92% of its atoms at the surface compared to 63% for a 147-atom cluster (or third complete shell), and to

45% for a 561-atom cluster consisting of 5-complete shells.81 The high percentage of

surface atoms in nanostructures results in a higher contribution of these atoms to

the physical and chemical properties of the material compared to their bulk structure

counterparts. A characteristic of surface atoms is the broken symmetry (i.e., dangling

bonds) which is believed to play an important role in new physical properties observed

at the nanoscale. Both nanoparticles and thin films have a high percentage of surface

atoms and, therefore, both can be studied to help understand possible surface effects

on magnetic properties.

The long-term outcomes of this fundamental and novel research has implications

in three major potential areas:

• Catalysts: several industrial processes utilize noble metal catalysts

usually in non-bulk reduced dimensions to maximize surface area. As

with any other consumables in industrial processes, catalysts could be-

come costly if they are not recycled and reused for multiple production

cycles. Ideally, studying and understanding the induction of magnetic

properties in metal nanostructures could lead to the manufacturing of

metal catalysts holding both properties of catalysis and magnetism. A

magnetic metal catalyst could be easily collected using a magnet, in

turn reducing cost. For example, Au nanoparticles show high catalytic

activity towards CO oxidation when nanoparticles are so small as to

40 be in a non-metallic state.82 In addition, studies have shown that con-

version reactions are accelerated when magnetic states are present at

the surface of metal catalysts,83 while others studies suggest that cat-

alytic activity of common catalysts could be matched by magnetized

surfaces.84

• Theranostics nanomedicine: the aim of theranostics (i.e.,acom-

bination of therapeutics and diagnostics branches of medicine) is to

simultaneously exploit nanosystems as diagnostics, drug delivery, and

monitoring therapeutics tools. 85 For example, iron oxide nanoparticles

have been utilized as diagnostic imaging agents (by exploiting their

magnetic properties) which could be upgraded to theranostic agents by

implementing therapeutic functions such as a drug delivery vehicle.85

• Biosensors: Nanoscale magnetic materials are useful in biosensing

due to their magnetic properties which are not found in biological sys-

tems.86 Ag and Au are biocompatible metals which is a highly sought

property for in vivo studies. Nanoparticles of these metals functional-

ized at the surface with specific biomolecules would be able to interact

with a target recognition site. Incorporating magnetic properties to

these nanoparticles could lead to an easy collection of the metal-target

complex with a magnet during retrieving procedures. A major advan-

tage of this system is the ability to concentrate target proteins that are

present at low titers.

1.8.1. Importance of magnetism

In general, the importance of magnetic materials is the vital role they have played in the development of civilization as a whole, starting with the loadstone compass

41 utilized in early maritime navigation,2 and continuing to be considered indispensable

in the technological development of modern society.5 For instance, at homes there are

several devices in which magnetism plays a central role in their operation including

electronics containing an electric motor.1 Other major technological fields influenced

by magnetism include medicine (e.g., magnetic resonance imaging),87 transportation

(e.g., magnetic levitation vehicle),88 information storage (e.g., hard drives in com-

puters),89 and the generation of electricity (e.g., hydroelectric power stations and transformers).90

Historically, the exploitation of magnetic materials has been a bulk property;

however, the field of nanoscience has provided fertile ground for a paradigm shift in

magnetism. Specifically, there is potential for size dependence to magnetic properties

of a material. This paradigm shift has moved the science of magnetism from the study

of bulk materials to nanodimensioned particles, wires and thin films. 2 It has been well

established that sizing plays an important role in the electronic, chemical and optical

properties of nanomaterials due to their higher ratio of surface-to-core atoms exhibited

in nanomaterials compared to their bulk counterparts, and to the confinement of

electrons versus their delocalization in bulk materials.36 These novel properties have

the potential to yield new and surprising magnetic properties in materials.

1.9. Outline of thesis

The structure of this thesis will be presented as follows. The experimental portion of

this work is presented in Chapter 2. The main message of this chapter consists of

the experimental parameters to prepare metallic nanoparticles, and our unorthodox

approach to collect thin films using a commercial prototype nanogenerator instru-

ment.

Chapter 3 will deal with an early approach to study our nanostructures utilizing

42 light as a probe to magnetic behaviour. Experimental work was performed with a

Magneto Optical Kerr Effect instrument specifically built as part of this thesis work.

The data acquired during this period set the tone for the following chapters.

Peer-reviewed data describing the physical characterization and magnetic be-

haviour of coinage metals Cu and Ag nanomaterials will be presented in Chap-

ter 4 and Chapter 5 respectively. These two chapters are published or ready-to- be-published work, which are presented as individual units with additional informa- tion and figures. The magnetic characterization was carried out with a commercial

Superconducting Quantum Interference Device (SQUID) magnetometer. The nanos- tructures studied in this work include nanoparticles and thin films synthesized by sputtering in the gas phase. Also, preliminary findings with Au nanostructures are presented in Chapter 6.

Finally, Chapter 7 will present a general overview and future work.

43 Chapter 2

Experimental details

2.1. Nanomaterial synthesis in the gas phase

The equipment utilized for the preparation for our nanomaterials was a NanoGen sputtering apparatus developed by Mantis Deposition Ltd. This instrument was designed to generate gas-phase nanoparticles (NPs) under vacuum conditions with pristine surfaces (i.e., free of capping chemicals necessary in other synthetic methods).

The main advantage of this process is that it allows a direct comparison of bare surface nanomaterials versus those with functionalized surfaces, while maintaining size and metal composition effects as independent factors. All three factors (i.e., surface state, nanodimensionality, and material composition) seemed to have a direct or indirect effect on the magnetic behaviour experimentally observed and theoretically predicted in nanomaterials. Although our initial vision consisted of working exclusively with

NPs, we discovered an unorthodox procedure utilizing the same instrument to prepare thin films composed of micro- and nanostructures as experimentally demonstrated by microscopy imaging techniques.

44 2.1.1. Nanoparticle preparation

The nanogenerator consists of a three-chamber vacuum system. A photograph of the instrument is presented in Figure 2.1. In the aggregation chamber (Figure 2.1AC), a dc magnetron head sputters a metal target generating an atomic vapour, followed by condensation of these atoms in the presence of an inert gas cooling. Once formed,

NPs are swept out of the aggregation chamber into a quadrupole mass spectrometer

(Figure 2.1MS), and finally collected onto a substrate at the deposition chamber

(Figure 2.1DC). This configuration was used in order to determine all the parameters required for specific sizes of each metal used in these studies.

Figure 2.1: A photograph of a version of the Mantis NanoGen deposition system. The instrument consists of three components: (AC) the aggregation chamber which houses the metal target from which nanoparticles are generated, (MS) the mass filter which allows for in situ mass spectrometry, and (DC) the deposition chamber where the selected masses are collected onto a substrate.

The configuration shown in Figure 2.1 is not suitable for the preparation of densely packed NP samples. It is estimated that the mass filter placement in this particular configuration produces a particle flux reduction of 1-2 orders of magnitude at the deposition chamber as per our current measurements. For instance, preparation of a

45 high density monolayer sample takes approximately 10 min in a configuration without

the mass quadrupole (i.e., consisting only of the aggregation zone and the deposition

chamber) but it would take 100-1000 min in the setup presented in Figure 2.1 with

all three components assembled. This latter scenario was not feasible for our needs;

thus, the chamber configuration was changed accordingly. An early adoption setup

consisted of placing the mass spectrometer after the deposition chamber. The final

configuration assembly consisted of removing the mass spectrometer altogether, while

directly connecting the aggregation chamber to the deposition chamber in order to

deposit higher density samples. It is important to specify that bypassing the mass

spectrometer does not affect the resultant NP size, as the gas-phase nucleation and

growth of NPs is terminated by the time the NPs exit the aggregation chamber.91

NP diameter d is calculated based on the frequency f data from the mass spec-

trometer as follows:

f × amu ÷ ρmetal = VNP (2.1)

where amu is atomic mass units in kg and ρmetal is the density of the sputtered metal

3 3 in kg/cm . The resulting volume V NP in cm can then be substituted into the volume of a sphere equation in order to obtain the radius r. Although NPs are not perfect spheres, it is an appropriate model. NP diameter d is obtained by multiplying r by

2.

In the gas-phase condensation process, the NP diameter is controlled by three key factors: the aggregation zone length, the gas flow rate, and the material concentration within the aggregation zone – all of these affect, directly or indirectly, the residency time of the NPs in the aggregation zone. Each key factor is explained below:

• Aggregation zone length: The aggregation zone length can be manu-

ally adjusted via a linear translation arm varying the sputtering head

46 position.92 This length affects the residence time of the metal atomic

species by extending (reducing) the time it takes to exit the aggregation

zone, while in the process increasing (decreasing) collisions with other

particles.93 Essentially, the tuning of this multi-collision path length

manages NP sizes. For instance, the longer the aggregation zone length,

the more collisions (or probability of collisions) the atomic species ex-

perience during particle growth, resulting in larger particles as can be

seen in Figure 2.2.

Figure 2.2: The aggregation zone of the nanogenerator. (a) The main components of the sputtering head are magnet (red) encased in a housing serving as a cathode (green) over which the metal target sits. During sputtering, Ar+ ions (blue spheres) present in the plasma are attracted to the negatively-charged cathode located behind the metal target, while in the process striking the metal target primarily within the confines of a well defined circle shaped by the magnet. The constant bombardment of ions removes metallic atomic species from which nanoparticles (copper spheres) are formed. (b)Afreshand(c) used metal Cu targets. The circular groove in the front face of the metal target in c is a typical characteristic of plasma etching.

47 • Gas flow rate: Two gasses were introduced into the system during

sputtering: the carrier and the plasma generating gas – He and Ar,

respectively. The function of the inert carrier gas is to sweep the parti-

cles from the aggregation zone to, for example, the deposition chamber.

By increasing gas flow rates, the particles sweep through the aggrega-

tion zone faster, thus reducing the time for particle growth. In con-

trast, higher Ar flow rates present during sputtering also increases the

strength of the plasma.

• Material concentration: Particle growth is also affected by the concen-

tration of atomic species being sputtered from the metal source. This

is affected by two parameters: Ar concentration and current applied

to the dc magnetron. The presence of more Ar atoms and of higher

currents, increases the intensity of the plasma resulting in more metal

atoms ejected from the target. The increase of metal atoms concentra-

tion leads to more collisions resulting in bigger particles.

Figure 2.3 shows the general effects of these key factors during testing conditions for the generation of NPs. These conditions include the aggregation zone length in relation to the position of the target, the gas flow rate as a combination of Ar and

He flow rates, and the effect of concentration affected by the current applied during plasma generation. The corresponding parameters for each key factor used to generate each panel on Figure 2.3 are found on Table 2.1. Figure 2.3a shows the pronounced effect that the aggregation zone length has on NP diameter. This factor shows the most clear response on the sizing outcome of NPs. All three distance trials show the presence of peak with the shorter (longer) aggregation zone length producing the smaller (bigger) NPs as expected. Figure 2.3b depicts the complex effect of the gas flow during plasma generation. As can be seen in the figure, two sample

48 Sample NP diameter Ar flow rate He flow rate IVD (nm) (sccm) (sccm) (mA) (V) (cm) Figure 2.3a 5 cm 2.92 50.6 11.1 308 282.3 5 7 cm 3.68 50.6 11.1 308 282.3 7 9.6 cm 5.36 50.6 11.1 308 282.3 9.6 Figure 2.3b 70 sccm – 60.2 10.3 253 259.5 6 70 sccm 3.00/5.62 70.3 0.0 253 262.9 6 100 sccm 4.70 70.2 30.3 253 263.9 6 Figure 2.3c 254 mA 3.26 50.1 10.1 254 277.3 4 308 mA 3.38 50.1 10.1 308 286.1 4 335 mA 3.20 50.1 10.1 335 286.4 4

Table 2.1: Conditions for the generation of Cu NPs of different sizes. I is the sput- tering current, V is the sputtering voltage, D is the distance of the metal target to the aggregation chamber exit. trials were produced using the same total flow rate set to 70 sccm (standard cubic centimeter per minute). The difference between them was the presence or absence of

He gas. Based on these two trial, the presence of He gas seemed to adversely affect the synthesis of NPs as there is no clear peak, while the absence of He results in two peaks indicating the presence of two NP populations. However, the trial consisting of 100 sccm gas flow, as a combination of 70 sccm Ar and 30 sccm He, demonstrate a better set of parameters to synthesize NPs under these conditions. Figure 2.3c represents the current effect applied during plasma generation. From our data, it seemed that current has almost no effect on NP sizing, as the difference between all three diameters is almost unnoticeable. The parameter itself did not allow for a large testing window as a result of a minimum current value of approximately 250 mA to start the plasma, and, as shown in Figure 2.3c, higher values provided no benefit at least on this particular set of conditions.

This NP generation method allows for a very controlled size particle and rela- tively narrow size distribution93 as shown by our work.94,95 By varying the process

49 Figure 2.3: Effects of different key factors during the generation of gas-phase metal NPs. a The position of the metal target in relation to the NP exit orifice in the aggregation zone represents the aggregation zone length. b The gas flow rate as a combination of Ar and He flow also affects the residence time of the forming NPs. c The effect of material concentration affected by the current applied during plasma generation. All signals have been normalized for ease of visualization. parameters, size range can be tuned to sub-12 nm particles as determined by in situ mass spectrometry measurements. Additionally, NP diameters can be determined post-deposition via microscopy surface characterization techniques.

Another advantage of this preparation method is control over particle density deposited onto the substrates. The relative density of the NPs films is the result of deposition time. Higher density samples require longer deposition times than lower density samples. Moreover, inherently to the gas generation process, formed NPs carry a negative charge during preparation. As NPs are generated in the aggregation

50 chamber, they travel away from the source due to a combination of electromotive force and pressure differential. The metal target and the freshly formed NPs are both negatively charged resulting in a large Coulomb repulsion that drives the NPs out of the plasma region of the apparatus. This charge plays a factor in the spacing of the

NPs. Coulomb repulsion among the NPs forces a relatively constant distance (relative to the desired deposition time) when they are deposited onto a substrate.96 Together these characteristics further describe another advantage of the gas generation process: size and particle density are independent of each other.

2.1.2. Thin film preparation

The unorthodox preparation of thin films with this instrument was the result of an observation made during attempts to synthesize Cu NP of different sizes. As can be seen in Figure 2.4 the used Cu targets exhibit very vivid colouration.Since the entire system is under vacuum and no other materials or molecules are present during sputtering, the unexpected coloration was indicative of metal nanostructures redepositing back onto the target during the different conditions attempted for NP synthesis. The different colours observed in Figure 2.4 are the result of different

NP concentration and their accompanying surface plasmon resonance. ON-target deposition was puzzling since the metal target experiences high temperatures during sputtering as a result of the plasma constantly striking the target. In addition, the negative charge of both the cathode (located behind the metal target) and the NPs, nanomaterial deposition onto the target was not logical. Based on this observation, three possible theories were developed attempting to explain this phenomenon.

The possible scenarios for our observed self-depositing nanostructures are pic- torially presented in Figure 2.5. First, Figure 2.5a represents an scenario whereby formed NPs are positively charged and will re-deposit onto the metal target due to their Coulomb attraction to the cathode. This is similar to the Ar+ cations being

51 Figure 2.4: Three different used Cu targets showing vivid colours as a result of nanomaterials redeposition during NP synthesis. This observation was made during attempts to synthesize metal NPs of different sizes. attracted to the cathode during sputtering. This possibility, however, was quickly abandoned due to the fact that there is a high electron concentration at the aggre- gation zone as a result of the ionization process of Ar atoms to form plasma. This is substantiated by the fact that NPs created by this or similar methods are nega- tively charged, and thus are able to be separated by mass-to-charge ratios using mass spectrometry.

The second scenario depicted in Figure 2.5b consists of magnetic clusters formed during aggregation. Our results indicated that a small population of metal particles or species were depositing back onto the target as a result of the magnetic field ex- erted by the dc magnetron performing a self-filtering effect preferentially attracting magnetic particles. At the time these experiments were being performed, a freshly published manuscript claimed that synthesis of magnetic coinage and noble metal

NPs was possible by solution synthesis under a magnetic field. 97 Experimentally, the claim simply stated that performing solution synthesis of NPs in a beaker over a stirring plate (containing a magnet) led to ferromagnetic nanomaterials of Cu, Ag, and Au among other metals. Additionally, a different research group studying Ag

NP formation in solution discovered that the most popular pre-nanoparticle cluster

52 Figure 2.5: Possible scenarios depicting self-deposition of nanomaterials onto the metal target. (a) Scenario portraying the case of Coulomb attraction by positively charged metal NPs. In this case, the NPs are following the same path as the Ar+ in the normal process of sputtering. (b) Scenario in which the formed NPs are “inherently magnetic” and are attracted to the back to the target as a result of the magnet positioned behind the metal target required for sputtering. (c) Scenario depicting the case where outgoing metal nanomaterials experience collisions with incoming Ar+ ions. In this latter physical process, nanomaterials loose their potential energy required for traveling away from the target. See text for further details.

98 formed was the icosahedral Ag13. It was discovered that the Ag13 cluster is promptly formed in large concentrations before aggregating to form bigger NPs. Finally, theo-

retical work by a third separate research group hinted at the Ag13 cluster being the

most magnetic with a 5 μB magnetic moment per particle amongst Ag clusters of

2-22 atoms.99 This cumulative literature information led us to believe our system was

synthesizing 13-atom metal clusters which were being magnetically attracted towards

the metal target as a result of the magnet located behind the target and required for

sputtering. After numerous experimental tests, incoherent data, and a lack of hard

evidence, the literature misdirection became a cautionary tale of not to go down the

rabbit hole, and of being more critical when researching published material. Inter-

estingly, years later, the same research group claiming the synthesis of ferromagnetic

coinage and noble metal nanoparticles under a magnetic field, published a correction

citing Fe parasitic contamination as their source of ferromagnetic behaviour. 100

The third scenario presented in Figure 2.5c is believed to be the correct process for

53 the backscattering observed during our synthetic method. The process of sputtering ejects material that ideally interacts with other source material resulting in forma- tion and growth of NPs which are collected outside the aggregation chamber. The other possibility is that the ejected material interacts with its surroundings, such as target-inbound Ar+ ions, leading to backscattering. Redeposition is experienced by some of the sputtered material when multiple collisions with Ar+ ions a short distance from the metal target result in a loss of potential energy and their inability to travel farther, leading to backscattering.94,101,102 Additionally, collisions during backscatter- ing could render the NP neutral by transferring its electron to the striking species.

Thin films produced in this matter can be collected simultaneously with NPs or can be created under the same conditions as NPs, yet they are topographically different from NPs. Thin films do not contain discrete particles separated by some distance, as in the case for NP films, and are composed of a continuous metal surface. Thus the manufacturing process allows for a direct comparison between surface-free NPs and surface-free thin films produced under similar conditions.

For these experiments, He and Ar compressed gases were obtained from Praxair at 5.0 Ultra High Purity (equivalent to 99.999% pure). Gas flow rates were regulated by Mass-Flo controllers (MKS Instruments Ltd., Type 1179A, SEMI Gas Code 001 for He and 004 for Ar) adjusted using a MFC Controller (Mantis Corportation Ltd.).

The Ar plasma was ignited using a high voltage unit (TDK Lambda GEN600-1.3).

In order to achieve high vacuum during deposition, three pumps were connected to the NanoGen system: a Varian DS 302 dual stage rotary vane vacuum pump

(Agilent Technologies) and two Varian Turbo-V 550 turbomolecular high vacuum pumps (Agilent Technologies). Pressure inside the NanoGen was measured using a

Varian XGS-600 vacuum gauge controller (Agilent Technologies).

54 2.2. Optical setup

A photograph of the experimental setup is presented in Figure 2.6. An schematic

representation of the experimental optical setup is presented in Figure 3.8 found in

Chapter 3. The optical setup was mounted on a home-made breadboard made of

acrylic glass. The light source was a deuterium white lamp (Mikropak DH-2000).

Initially, light traveled through fiber optic before encountering the calcite Wollas-

ton prism (Edmund, #68-821) or first polarizer. Incident light beam (i.e.,either

p-polarized, s-polarized, or non-polarized light) was reflected from the sample to- wards the detector. Reflected light first encounter a lens allowing for a better signal collection, then a second polarizer, and finally the detector (Ocean Optics USB4000

Fiber Optic Spectrometer). The optical setup was dismantled at the end of these studies to use the same components for the magneto-optical Kerr effect setup

2.3. Magneto-optical Kerr effect setup

A photograph of the experimental setup is presented in Figure 2.7. An schematic representation of the setup for measuring polar Kerr rotation during magneto-optical

Kerr effect (MOKE) experiments is shown in Figure 3.6 found in Chapter 3.

Components of the MOKE setup were mounted on a home-made breadboard made of acrylic glass and a Al metal sheet. The light source varied during experimentation starting with a deuterium while lamp (Mikropak DH-2000, data not shown) and set- tling on a 532 nm YAG green laser (Lotis TII) in order to monitor a single wavelength.

The light intensity was followed and measured using an Optic Ocean spectrometer

(USB4000). Light was polarized by a calcite Wollaston prism (Edmund, #68-821,

P1 on Figure 2.7) in either the p-ors-polarization before probing the sample. All tested samples were deposited on NaCl prisms and probed either as internal or ex- ternal reflection. Reflected light passed through a lens in order to concentrate the

55 Figure 2.6: A photograph of a version of the experimental optical setup. Light produced by the light source (LS) travels through optic fiber before encountering the first polarizer (P1). Polarized light probes the sample (S) and after reflection, it is focused using a lens (S), before encountering the second polarizer (P2) and finally reaching the detector (D). The sample in the picture is set for internal reflection.

reflected light. Reflected light then pass through a second polarizer (polarization filter

80 #G036325000 from Qioptic, P2 on Figure 2.7) before reaching the detector.

The magnetic field was generated with a C-frame electromagnet (GMW Asso-

ciates, 45 mm #3470) using a bipolar operational power supply (Kepco, BOP50-8M-

4886). The electromagnet was calibrated using a LakeShore gaussmeter (475DSP

model, 475-HMNT-4E04-VR) equipped with a Hall probe (HMNT-4E04-VR). Elec-

tromagnet calibration is shown in Figure 2.8. The maximum current allowed by the

electromagnet without cooling was ±3.5 Amp. Current values were adjusted manually during both calibration and experimentation.

56 Figure 2.7: A photograph of a version of the experimental MOKE setup. Light produced by the light source (LS) travels through optic fiber before encountering the first polarizer (P1). Polarized light probes the sample (S) under a magnetic field (M). Light reflected from the sample is focused using a lens (S), before encountering the second polarizer (P2) and finally reaching the detector (D). Magnet poles were covered in black rubber as a measure to avoid stray light reflecting from their metal surface.

2.4. Superconducting quantum interference device magnetometry

The technique for analyzing magnetic materials with utmost precision is the super- conducting quantum interference device (SQUID) magnetometry103 with theoretical lower sensitivity limits of approximately 10−7 emu.4 A drawback for reaching such low sensitivity is the lengthy measurements as a results of the dynamic magnetic field values applied, for example, during M vs μ0H loop measurements. In some cases these measurements could take approximately 12 hours depending on the sequence parameters.

All measurements were performed on a field-shielded Quantum Design MPMS XL-

7S SQUID magnetometer system. Metal NPs and thin films were tested by mechan-

57 Figure 2.8: Calibration curve for the electromagnet utilized during MOKE experi- ments. Maximum current in our setup was ±3.5 Amp. Resulting magnetic field was monitored with a Hall probe. ically exfoliating them with Kapton tape from their substrates (either Si or NaCl).

Alternatively, in some instances, nanomaterials were tested without exfoliation (i.e., measurements gathered together with the substrate they were collected on). For ei- ther approach, samples were subsequently inserted into a clear diamagnetic plastic straw for measurement. Isothermal measurements of magnetization, M, as a function of field strength, μ0H, were carried out at 300 and 5 K (latter data not shown) by cycling the applied magnetic field strength between 1 and –1 T.

In general, once the analyte is loaded into the instrument, SQUID magnetometry

measurements are performed by an up-and-down sample motion through current-

carrying superconducting detection “pickup” coils as depicted in Figure 2.9. Since the

pickup coils are constructed from superconducting material, a constant current flow

without loss is achieved once current is injected into them. As previously described,

current flowing in a wire creates magnetic flux lines outside its physical boundaries.

The pickup coils are coupled to a small SQUID unit loop, which is in itself a tiny-

58 sized component of the entire SQUID magnetometer apparatus. The key role of the

SQUID unit loop is as a magnetic flux detector.

A conventional SQUID loop consists of a small ring of superconducting material

cut in half by two parallel non-superconducting barriers103 known as Josephson junc-

tions. Wire leads are connected at each end of the ring allowing for current to pass

through it.103 As input, current is injected at one end of the ring and removed at the other; whereas the output is the voltage that develops across the gap. 104 Varia-

tions induced in the current of the pickup coils circuitry produced by the magnetic

field lines of a magnetic sample leads to a current flow in order to compensate the

flux change.15 These current variations are in turn translated to variations in the

SQUID output voltage.105 Thus, in SQUID magnetometry measurements, there is no

measurement of magnetic moments but of magnetic flux interfering with the mon-

itoring voltage. Additionally, the SQUID loop sensor does not produce an output

voltage for a uniform magnetic field or a constant field gradient.106 Therefore, when

a uniform structure longer than the span of the pickup coils, such as the diamagnetic

plastic straws utilized to load our samples, produces a constant flux which does not

contribute to a magnetic signal.106

The Josephson junctions present in the SQUID loop sensor are approximately 1

nm in length104 and act as physical barriers for the electrons which must quantum

tunnel through them in order to reach the output half of the SQUID loop. Quantum

tunnel, or most commonly referred to as tunneling, is a phenomenon experienced by

subatomic particles when transitioning between states that are separated by a poten-

tial barrier.9 Quantum objects posses a wave nature which allows them to transition

between two states without climbing the potential barrier, instead “tunneling” to

reach the target state.107 Because the current passing through the SQUID unit can

take either path around the ring, electron pairs acting as quantum mechanical wave-

59 Figure 2.9: Scheme of SQUID magnetometry apparatus showing its major compo- nents. During magnetometry measurements, a sample experiences up-and-down mo- tion (double-headed black arrow) in the presence of a magnetic field. The magnetic field lines of the sample interfere with the current carried by the superconducting wire and pickup coils. The change in current flow is detected by the SQUID loop sensor. Red gaps in the SQUID loop sensor represent Josephson junctions. Blue single-headed arrows show the direction of the current. Adapted from Ref. 3 functions crossing the junctions can constructively interfere (in-phase) or destructively interfere (out-of-phase) when they meet on the other side of the structure.103,104 The way electron wave functions align is set by the applied magnetic field acting on one side of the SQUID loop.103 This occurs because the magnetic field changes the phases of the quantum wavefunctions across the junctions.108 Thus, the current sensitivity through the ring depends on the type of interference created by the interaction of the magnetic field with the SQUID loop103 turning the SQUID sensor into a linear current-to-voltage converter.105

Diamagnetic corrections were performed in all our M vs μ0H loop measurements. These corrections were made in order to account for the magnetic susceptibility of the Si substrate (i.e., for samples without exfoliation), the Kapton tape, and the non-surface bulk interior of our metal nanomaterials samples, all of which behave diamagnetically.

60 2.5. Magnetic quartz crystal microbalance

Magnetic quartz crystal microbalance (MQCM) is a recent application described by

Janata and co-workers109,110 and it is derived from the mature and robust quartz crystal microbalance (QCM) technique. The essence of the MQCM technique is its ability to qualitatively differentiate the magnetic properties of materials.

In a regular QCM experiment, a thin quartz crystal consisting of two electrodes deposited on opposites sides109 oscillates at a nominal frequency (in our case 5 MHz).

The thin crystal is a piezoelectric material that undergoes shear deformation (i.e., resonates at the nominal frequency oscillation)111 when an alternating-current volt- age is applied across the opposing electrodes generating a transversal acoustic wave propagating through the crystal. 112 Mass changes (either positive or negative mass changes) on the face of the crystal affect the propagation of the resonant frequency of oscillation in turn indicating a response by the resultant frequency change. Thus,

QCM is a well understood technique used as a mass-sensitive detection system able to measure, in some cases, picogram amounts of materials. As a result of its sensitive nature, the idea behind MQCM is to be able to perturb the frequency of oscillation of the quartz crystal by depositing magnetic material onto it, followed by the appli- cation of a magnetic field, and measuring the frequency change of the system pre- and post-magnetic field exposure.109,110 For a fixed mass of material present on the quartz crystal face, in the presence of a magnetic field, magnetic material, for exam- ple, experiences attractive forces that also affect the frequency at which the QCM oscillates. It is then argued that the change in QCM oscillation frequency observed in the presence of the magnetic field is directly related to the strength of the interaction of the analyzed material with the magnetic field.

A schematic of the major components of the MQCM setup utilized to study mag- netic properties is shown in Figure 2.10. The setup consisted of a QCM crystal and

61 a bar magnet. For the initial proof-of-concept MQCM experiment presented in Fig- ure 4.5b found in Chapter 4, the distance between the magnet and face of the quartz crystal was varied. The magnitude of the magnetic field applied to the quartz crystal was based on separate measurements of the distance between a Hall probe and the bar magnet. The Hall probe is a gadget magnetic field sensor consisting of a flat tip where the sensor is located and it is utilized to measure magnetic field strengths. Its sensor capabilities is based on the Hall effect. The Hall effect is measured when a current-carrying conductor is exposed to a magnetic field generating a potential dif- ferenceacrossit(i.e., the flat tip) in the direction perpendicular to both the current and the magnetic field.1

Several magnetic field strengths were applied based on distance to representative magnetic samples and their frequency change recorded. A QCM crystal with a thin

film of high vacuum grease covering one of its electrodes was considered the blank measurement. The grease was deposited prior to material deposition in order to fix the material in place to avoid loss during oscillation or exposure to the bar magnet. Using this approach, the magnetic responses of ferromagnetic Fe, weakly paramagnetic Al, and the diamagnetic blank quartz crystal were measured semi-quantitatively. The Fe and Al samples consisted of filings collected from bulk materials, and thus considered non-nanodimensioned.

As can be seen in Figure 4.5b, the QCM crystal coated with a film of high vac- uum grease shows a frequency change with increasing magnetic field strength. The effect is relatively minor and the decrease in frequency is opposite to the increase in frequency observed when ferromagnetic material is present on the QCM face. When ferromagnetic Fe filings are embedded in the vacuum grease, a much larger frequency change increase is observed. The increase is several factors larger than the decrease observed without magnetic material present. For comparison purposes, Al filings (a

62 Figure 2.10: MQCM concept to study magnetic properties. (a)FrontviewofaQCM crystal. (b) Side view of the setup including the quartz crystal and the bar magnet (red cylinder) moving several d distances to exert different magnetic field strengths (green arrow).

weak paramagnet) is also plotted in Figure 4.5b. Trend comparison of the Al filings

with the blank, show the presence of Al filings has no obvious additional effect on

the QCM oscillation frequency. Finally, the semi-quantitatively nature of these data

cannot determine magnetic moments of the analytes; however, the trends presented

here are a representation of the behaviours of different magnetic materials during an

MQCM experiment.

2.5.1. In situ MQCM

Although SQUID magnetometry is recognized as definitive in determining magnetic

properties, one of the experimental challenges with preparing samples under a vacuum

inside our NanoGen apparatus was to retrieve them for further analysis and, while in

transit, maintain their surface in a pristine state. Accordingly, although NP samples

and thin films were prepared under vacuum conditions and kept under either Ar

or O2 gas, the reactive nature of NPs combined with the unavoidable exposure to ambient air during sample transfer (i.e., during sample retrieval after initial synthesis

or during SQUID sample preparation and loading) could lead to partial oxidation of

the NPs, even those stored under Ar gas.

63 To study the effect of inevitable surface oxidation, the magnetic properties of pris-

tine NPs were also studied in situ under vacuum conditions using MQCM technique.

MQCM measurements were acquired by initially depositing metal NPs onto a fresh

QCM crystal and monitoring the change in frequency during magnetic field exposure.

Figure 2.11 shows the implementation of the MQCM technique inside our apparatus

for in situ measurements. For MQCM experiments, QCM crystals with a frequency of 5.000 MHz (Lap-Tech Inc.) were loaded onto a quartz crystal microbalance ana- log controller (Stanford Research System, model QCM 100) and measurements were recorded with a universal counter (Agilent, model 53131A). Magnetic fields were measured using a Hall probe (LakeShore, HMNT-4E04-VR model) connected to a

LakeShore 475 DSP gaussmeter. The resulting NPs and crystal were immediately exposed to inhomogeneous magnetic field strength of 0.1 T from a bar magnet. Thus, the change in frequency (Δf MQCM in an MQCM experiment) is defined as the differ- ence between the frequencies in the presence and in the absence of a magnetic field

(ZF), that is, Δf MQCM = f0.1T – fZF. This experiment provided information about the magnetic properties of the pristine metal nanomaterials. Then, the vacuum in the

NanoGen instrument was broken when ambient air was slowly leaked into the system,

and once again, MQCM measurements were performed. Further oxidation of these

NPs was achieved by mild heat treatments ex situ followed by MQCM measurements.

The MQCM technique has the added advantage of allowing for precise monitoring

of the mass of oxygen uptake by the metal NPs as a result of exposure to air. To

further explore the effect, more extensive oxidation was driven by incubating thin

films in air at mild temperatures (e.g., ≤100 ◦C) and then analyzing them by SQUID

magnetometry. All MQCM measurements were performed at room temperature.

64 Figure 2.11: MQCM setup to study magnetic properties. (a) Back and (b)frontview of a QCM crystal set for an in situ MQCM experiment. Diamagnetic materials were used in order to avoid interactions with the bar magnet. For example, a plastic shell was used as a holder and Cu alligator clips were used as electrical contacts. As can be seen in (b) an orifice through the plastic shell allows for NP deposition exclusively onto the electrode of the QCM crystal. (c) Shows the setup inside the collection chamber. In this instance, the bar magnet is position in an ON position but it can be manually retracted to exert virtually no magnetic field strength (i.e., OFF position) to the sample. Again, to avoid interactions with the bar magnet, the wires connecting the system to the remaining controller and counter located outside the chamber were positioned away from the bar magnets.

65 2.6. X-ray photoelectron spectroscopy

Oxidation states and sample purities were tested via X-ray photoelectron spectroscopy

(XPS). A Physical Electronics PHI VersaProbe 5000-XPS instrument was used to

record the XPS spectra. The spectra were collected using a monochromatic Al Kα

source (1486.6 eV) at 50 W, and with a beam diameter of 200.0 μm. Some sam- ples were further analyzed in situ after Ar sputtering at a rate of 10 nm/min for 5 min to reveal a fresh surface. All spectral analyses were performed with CasaXPS software.113 Spectra were corrected by calibrating all peaks to the adventitious C

1s signal at 284.8 eV. A Shirley-type background was used, and curve fitting was

performed using a combination of Gaussian and Lorentzian [GL(85)] profiles. The

analysis of the O 1s region followed a methodology presented elsewhere.114

2.7. Atomic force microscopy

Atomic force microscopy (AFM) topography images were collected using a Veeco

Innova AFM instrument mounted on an acoustic-vibration isolation system. Images

were collected under ambient conditions in tapping mode using a silicon cantilever

with a spring constant of ∼50 N/m.

2.8. Scanning electron microscopy

Scanning electron microscopy (SEM) images were collected using a Zeiss Σigma VP

field emission scanning electron microscope equipped with an Oxford INCA X-Act unit for energy-dispersive X-ray spectroscopy (EDXS) analysis.

66 2.9. Powder X-ray diffraction

Powder X-ray diffraction (PXRD) was performed on a Bruker D8 Advanced (ECO) polycrystalline X-ray diffractometer with Cu Kα radiation from a fine focused sealed

X-ray tube (40 kV, 25 mA) and LYNXEYE XE 1D detector in continuous mode. The diffraction radius was 280 mm, divergence slit was set to 0.6◦, anti-scattering slit was widely open, and the positive sensitive detector (PSD) opening set to 2.2◦.Thescan was performed from 20 to 90◦ (2θ)in0.01◦ increments.

2.10. Experimental methodology for Cu nanomaterials

Pristine Cu NPs of high purity were fabricated using the gas-phase synthesis ap- proach with control over NP diameter in the range from 4.5 ± 1.0to9.0± 1.8 nm. All samples were prepared under vacuum using a Mantis NanoGen sputtering system as described elsewhere.96 The metal target was cooled with a closed 10 ◦C water loop system at all times. During standard use, as seen in Figure 2.12, NPs generated in the aggregation chamber are swept out to adjacent chambers as a re- sult of interchamber pressure differentials. Adjacent to the aggregation chamber is a quadrupole mass spectrometer, followed by a collection chamber where NPs are deposited onto substrates. As the metal NPs produced in the aggregation chamber are negatively charged,115 mass spectrometric methods were used to analyze the NPs produced. Thus, through mass analysis preparative conditions yielding specific size distributions of NPs were determined as described in Table 2.2. Note that, during NP collection (site a in Figure 2.12), the mass spectrometer chamber was removed, and the aggregation chamber connected directly to the collection chamber to maximize the NP yield. It is important to specify that bypassing the mass spectrometer cham- ber does not affect the resultant NP size, as the gas-phase nucleation and growth of

NPs is terminated by the time the NPs exit the aggregation chamber.91

67 NP diameter Ar flow rate He flow rate IVDp (nm) (sccm) (sccm) (mA) (V) (cm) (Pa) 4.5 ± 1.0 50.1 10.1 308 287.4 6 32.0 6.5 ± 1.0 50.8 10.4 308 283.5 9 32.0 9.0 ± 1.8 70.7 30 253 275.9 9 56.0

Table 2.2: Conditions for the generation of Cu NPs of different sizes. I is the sput- tering current, V is the sputtering voltage, D is the distance of the metal target to the aggregation chamber exit, p is the pressure in the aggregation chamber during deposition.

In addition to collecting NP samples in the deposition chamber, thin films of

Cu were also collected by placing a substrate directly onto the metal target away

from the plasma etching location (site b in Figure 2.12). The process of redeposition

is experienced by some of the sputtered material when multiple collisions with, for

example, gas atoms a short distance from the target result in a loss of energy and

their inability to travel farther, leading to backscattering.101,102 Both NPs and thin

films were collected on either Si(100) wafers (∼5 × 5mm2 in size) or NaCl crystals

2 (∼10 × 7mm in size). Following preparation, samples were stored under Ar or O2 until further analysis. Depositions were performed in 5-min intervals (i.e.,5minof

deposition followed by 5 min of rest) as an attempt to aid with the metal target cooling

as the kinetic energy of the Ar ions striking the surface leads to local heating. 116

2.11. Experimental methodology for Ag nanomaterials

All samples were synthesized under vacuum using a Mantis NanoGen sputtering sys-

tem94,96 and schematically illustrated in Figure 2.13.

Different size Ag NPs were prepared by modifying source conditions to generate

NPs with modal diameters of 3.3 ± 0.9, 5.0 ± 0.9, and 7.8 ± 1.3 nm. Table 2.3 displays the parameters utilized during the preparation of Ag NPs.

In addition to collecting NP samples in the collection chamber, thin films of Ag

68 Figure 2.12: Schematic of deposition system showing only its major components. In the aggregation chamber, the red holder represents the dc magnetron, on which the Cu metal target sits. Site a (b) refers to the NP (thin film) sampling location. See text for full description. were also collected at two different locations within the aggregation chamber. Thin

films were prepared by placing a substrate directly onto the metal target away from the plasma etching location (site b in Figure 2.13). At this position (referred to as

ON-target sample herein) Ag thin films form on the substrate in the presence of a

250 mT magnetic field. The magnetic field is due to the presence of a dc magnetron behind the target as required for sputtering. Ag thin films were also collected on substrates placed inside the aggregation chamber but away from the metal target (site c in Figure 2.13). At this position, metal atoms from the plume land onto substrates as in atomic vapour deposition. These OFF-target samples were collected in the absence of any significant magnetic field; thus, the two film types were synthesized

NP diameter Ar flow rate He flow rate IVDp (nm) (sccm) (sccm) (mA) (V) (cm) (Pa) 3.3 ± 0.9 50.5 80.7 253 266.3 6 61.3 5.0 ± 0.9 50.1 50.7 253 259.8 9 50.6 7.7 ± 1.3 90.2 0 235 263.1 9 45.3

Table 2.3: Conditions for the generation of Ag NPs of different diameters. I is the sputtering current, V is the sputtering voltage, D is the distance of the metal target to the aggregation chamber exit, p is the pressure in the aggregation chamber during deposition.

69 Figure 2.13: Schematic of deposition system showing only its major components. In the aggregation chamber, the red holder represents the dc magnetron, on which the Ag metal target sits. Sampling location are as follows: site a for NPs, site b for ON-target thin films, and site c for OFF-target thin films. See text for full description. under different conditions. Si wafers (∼5 × 5 mm in size) and NaCl crystals (∼10

× 7 mm in size) were used as substrates. Once synthesized, NPs and thin films were stored in sealed containers under either Ar or O2 atmospheres until analysis was performed. Finally, all depositions were performed in 5-min intervals (i.e.,5minof

deposition followed by 5 min of rest) to avoid overheating of the metal target.

70 Chapter 3

Optical studies of silver nanomaterials

3.1. Introduction

Light is an electromagnetic wave containing an electrical component and a magnetic

component perpendicular to each other, while both concurrently set to travel trans-

versely to the propagation direction.15 These two components allow for the indirect

manipulation of light with either an electric field or a magnetic field. In the present

study, the interaction of light with matter involves the electronic structure of the lat-

ter117 resulting in the ability to probe magnetic materials. The interaction between

light and matter when the latter is subjected to an external magnetic field H is called

magneto-optics.118 The main objective of this work is to study metal nanomateri-

als using magneto-optics to evaluate their magnetic behaviour (or lack thereof) as

a pre-selection process prior to more comprehensive magnetometry studies. There

are several magneto-optic phenomena which have been discovered by the scientific

community. These phenomena include the Zeeman, Voight, Faraday, and Kerr ef-

fects.15,118,119 In the following sections, only the latter two magneto-optic effects will be introduced with particular emphasis on the Kerr effect which was utilized in this

71 work.

3.2. Magneto-optical effects

In general, all magneto-optical phenomena study the interaction of light with a mate-

rial under magnetization. Specifically, with the exception of the Zeeman interaction,

all magneto-optical effects describe the polarization response of the probing light in-

duced by the interaction with a material’s magnetization M as a result of its own internal net magnetic moment μ (e.g., a ferromagnet)118 or by the effect of an exter-

nally applied magnetic field H.119

3.2.1. Faraday effect

Michael Faraday discovered the first magneto-optical effect that bears his name (i.e.,

the Faraday effect) in 1845.117 His discovery made the connection for the first time

between magnetism and light.15 The Faraday effect consists in the rotation of linearly

polarized light when the beam of light propagates through any transparent substance

(i.e., transmission, defined as the lack of interaction between light and material)

placed in a magnetic field,120 where the magnetic field H is parallel to the propagation

direction of light k,121 as seen in Figure 3.1. As previously mentioned, light has two

components which are perpendicular to each other; however, it is widely accepted

that when light is polarized, it is strictly referring to the alignment of the electric

122 wave component E c. This standardization is the result of the general effect of

light on matter involving the redistribution of charges, which an electric field has

a large effect on compared to a magnetic field.122 In essence, Faraday discovered

that when the electric component of plane-polarized light passes through glass in

the direction parallel to the applied magnetic field H, the plane of polarization is

rotated.119 Faraday’s original work was performed on transparent paramagnetic glass;

72 however, the effect is also observed in transparent diamagnetic, ferromagnetic, and ferrimagnetic materials.15 Additionally, a characteristic difference amongst materials is that diamagnetic specimens show a smaller rotation compared to materials with permanent magnetic moments.123

Figure 3.1: Graphical representation of the Faraday effect. The horizontal gray cylin- der is a transparent object through which the light is transmitted. Both the external magnetic field H (i.e., magnetization M ) and the propagation of the linearly-polar- ized light k vectors are parallel. The angle θF is the angle of rotation of the linearly polarized light caused by the Faraday effect and is referred to as the Faraday rotation. The rotation of θF has been grossly exaggerated as a visualization aid. Adapted from Ref. 15.

The Faraday phenomenon can be explained by considering plane-polarized light as a superposition of two oppositely rotating circularly polarized components,3 as shown in Figure 3.2a. When planed-polarized light enters a medium in the presence of a magnetic field, one of the circularized components propagates faster than the other due to the different magnetically-induced refractive indexes for each component; thus, this process induces birefringence in the material (i.e., two indices of refraction within the same material).3,119 Recall that the refractive index n is defined as the ratio of the speed of light in vacuum, c, to its speed in a particular medium, c:3

c n = (3.1) c As a consequence of the difference between the left circular-polarized component refractive index nL and the right circular-polarized component refractive index nR, the propagation velocities of the two circularly polarized waves, c/nL and c/nR,dif-

73 fers.119 Most importantly, when they leave the sample, the two circularized-polarized components are out of phase with each other and their superposition results in plane-

3 polarized light rotated through an angle θF compared to the incoming beam (Fig-

ure 3.2b), where the Faraday rotation θF angle is proportional to the difference in

3 refractive index, nR–nL.

Figure 3.2: (a) Graphical representation of the electrical component E c of linearly polarized light and its superposition of left and right circularly polarized light compo- nents. Observed from the point of view facing the oncoming light beam. (b) Faraday rotation when the speed of one of the circular polarized components is affected by the medium, in effect rotating the linearly polarized light by an angle θF . Adapted from Refs. 3 and 124.

Although the Faraday effect was originally discovered in transparent glass, the ef-

fect is also present in materials that partially absorb (i.e., not completely transparent)

some of the incident light.125 As will be described in the next section, in this instance

the linearly-polarized light is not only rotated but also becomes elliptically-polarized.

The importance of the discovery of the Faraday effect is two-fold in that a lon-

gitudinal magnetic field H made all substances appear optically active,120 and this

effect is responsible for the affirmation of the electromagnetic nature of light. 117

74 3.2.2. Kerr effect

In 1876, John Kerr discovered a similar effect to Faraday’s.117,121 In simple terms, the Faraday effect deals with rotation of transmitted light, while the Kerr effect deals with rotation of reflected light.121 Additionally, the reflected light is also elliptically polarized in comparison to the incident plane-polarized light.119 Thus, the Kerr effect is the result of rotation and ellipticity changes compared to the linearly-polarized incident light.126 The Kerr effect is presented pictorially in Figure 3.3. In terms of applicability, although the Kerr effect has been known for more than a century, experimental work on magneto-optics as a technology for data storage did not begin until the 1950-60s.127

Figure 3.3: Graphical representation of the Kerr effect: (a) when polarized light is absorbed the intensity of one of the circularized components is reduced resulting in ellipticity, and (b) when the speed of one of the circularized components in the medium changes the overall effect is the resulting reflected light has changed its ellipticity ηK and its rotation θK resulting in the Kerr angle φK . Adapted from Ref. 124.

The Kerr effect can be described in the context of the the Lorentz force128,129 and the microscopic quantum theory.126,130 These explanations are discussed next.

75 3.2.2.1. Lorentz force

The semi-classical physical explanation for the Kerr effect can be understood in terms

of the Lorentz force as a basic interaction between the light and the electrons on the

magnetized surface.131,132 For this explanation to be valid, it is assumed the electrons

are localized within the atoms of the analyte,129 as opposed to the delocalized “sea-

of-electrons” concept widely accepted for metallic materials. In a polar geometry,

in the absence of a magnetic field, as the linearly-polarized light interacts with the

electrons on the sample, the oscillating electrical field E c of the light wave induces

an acceleration (i.e, motion) of these electrons along the plane of light polarization

at the surface, for example wholly on the x-axis.128,130 However, in the presence of

an external magnetic field (which is parallel to the light plane of incidence in a polar

configuration), an additional force acts on the motion of electrons. 130 This is the

Lorentz force F which on a moving electron is defined as:

F = −e(E c + v ⊗ B ) (3.2)

where e is the charge of an electron, E c is the electric field, and v ⊗ B is the cross

product of the movement of the electron v and the magnetic flux density B with a mag- netization perpendicular to the sample’s surface. This induces a new y-component to the amplitude of the probing linearly-polarized light originally oscillating in the x-axis with the reflected light emerging elliptically-polarized.129

3.2.2.2. Microscopic quantum theory

On a microscopic level, the quantum theory explanation of the Kerr effect is based on

the the spin-orbit coupling119,133 and the exchange interaction128 phenomena. When

plane-polarized light interacts with a ferromagnet, for example, the symmetry of the

left- and right-hand circularly polarized light is broken resulting in a separation of the

two superimposed components.119,134 In this process the refractive indexes of the two

76 circularly-polarized components are affected resulting in the elliptically polarization

outcome of the Kerr effect.119

Spin-orbit coupling (SOC) is an intra-atomic magnetic interaction between the

spin magnetic moment of the electron μspin, and the magnetic field created by the the

electron’s own orbital motion about the nucleus – where the electron is in an orbital

other than s orbitals.1,3,18,20 In addition, as shown in Figure 1.10 on Chapter 1, due

to exchange interactions the spin up ↑ and spin down ↓ electron populations are

separated by different lower energy levels. The combination of these two phenomena

are believed to be responsible for the Kerr effect. This is depicted pictorially in

Figure 3.4.

When light is absorbed, there are certain selection rules in quantum mechanics

that only permit certain transitions while denying forbidden cases. In the transition

metal case depicted in Figure 3.4, the selection rule Δl allows transitions between the

d and the p orbitals.128 Overall, Δl allows ±1 transitions. Similarly, the selection rule Δml allows for ±1 transitions where Δl=+1 and Δl=–1 correspond to left- and right-circularly polarized light, respectively.128 By definition, the values of the

quantum number ml are ±l. The transitions from the spin-orbit splitting in the 3d orbitals to the empty 4p orbital are distinguished by the photon polarization135 stated

by the selection rule Δml. Circularly left polarized light is absorbed when Δml=–1,

128 while circularly right polarized light is absorbed when Δml=+1.

When both spin-orbit and exchange interactions are present in a material, the

absorption spectrum of left- and right-circularly polarized light is different as a result

of the electronic structure of the material being modified differently by each polar-

ization128 as shown in Figure 3.4d. The net effect on the incident linearly-polarized

light is the different effects on its circularly-polarized components due to the unequal

refractive indexes nR and nL affecting rotation, and extinction coefficients κR and κL

77 Figure 3.4: Simplified pictorial representation of an atomic energy scheme for a 3d transition metal with distinct spin-up and spin-down electron populations. There are four stages depicting the presence or absence of exchange interaction and spin-orbit coupling (SOC). (a and c) Absence of exchange interaction. (a and b) Absence of SOC. (d) Presence of exchange interaction and SOC. Blue and red arrows present possible absorption transitions occurring between the occupied 3d orbitals and the unoccupied 4p orbitals according to the selection rule Δml=±1. In the presence of exchange interactions, the orbitals are separated by E ex or exchange energy. Adapted from Ref. 128.

78 affecting absorption.

As closing remarks for this section, it can be said that historically the magneto-

optic effects were discovered in simple systems which could be relatively easily ex-

plained. In metallic systems, however, as the ones studied in this work, due to their

strong adsorption characteristics,130 explanations are difficult to establish.126 In both

phenomena, the Faraday and Kerr effects were discovered in nonmagnetic materials

resulting in at least rotation of polarized light136 with an ellipticity factor if adsorp-

tion of the incident light is part of the system. Phenomenologically for both effects, a

magnetic material (i.e., possessing an internal magnetic moment μ) in the absence of

an external magnetic field H behaves identical as a nonmagnetic material in the pres-

ence of H.136,137 Thus, these magneto-optical phenomena are observed in all materials including insulators, semiconductors and metals (e.g., diamagnets and paramagnets); but, these effects are strongest in magnetically ordered materials such as ferromag- nets and antiferromagnets.9 A major disadvantage of using visible light (400-700 nm

wavelength) is the typical spatial resolution of optical techniques is ∼0.2 μm(using the Rayleigh criterion) which is above the domain size of, for example, 0.1 to 1 mm for bulk antiferromagnetic NiO.138 Recently, with new technical developments, the

Faraday and Kerr effects have been studied with X-rays (0.01-10 nm wavelength).137

Furthermore, X-ray magneto-optical spectroscopies allow for probing selectively the

magnetic state of individual elements in a material.139

3.3. Magneto-optical Kerr effect

The magneto-optical Kerr effect (MOKE) is defined as the change in the intensity

and polarization of light during its reflection from the surface of a sample under a

magnetic field regardless of whether the magnetic field is external (H )orinternal

(μ).9 The net effect on the probing incident plane-polarized light is reflection as

79 elliptically-polarized with the major axis of the ellipse rotated with respect to the

incident light119 as seen in Figure 3.3.

An experimental MOKE apparatus can be set-up in several configurations based

on the applied external magnetic field H (more specifically, the magnetization vector

M on the sample), with respect to the plane of incidence of the linearly-polarized

light, and to the surface plane of the analyte.140 The three basic configurations are

illustrated in Figure 3.5. These geometries are believed to potentially affect the

Kerr effect as the film surface is probed differently in each instance. The MOKE

configurations are:9,119,140

• Polar: in which the magnetization M vector is normal (i.e., perpendic-

ular) to the reflective surface of the sample and parallel to the plane of

incidence of the light;

• Longitudinal (also called meridional): in which the magnetization M

vector is parallel to both the reflective surface of the sample and to the

plane of incidence; and,

• Transverse (also called equatorial): in which the magnetization M vec-

tor is parallel to the surface sample and is perpendicular to the plane

of incidence.

Although MOKE magnetometry is considered a powerful technique to study prop- erties of known-magnetic materials such as magnetic anisotropy, configuration of magnetic domains,4 and magnetization reversal,140,141 it is also utilized as a qual-

itative pre-screening diagnostic alternative132 to other more demanding and costly

magnetometry techniques. Thus, our set-up attempted to probe our metallic films

to determine whether or not they were “magnetic”, rather than quantify their mag-

netism. The scheme of our home-assembled MOKE set-up composed of commercially

80 Figure 3.5: Magneto-optical Kerr effect (MOKE) in (a) polar, (b) longitudinal, and (c) transverse configurations. Green arrows indicate the orientation of magnetization M due to the applied magnetic field on the blue-slab substrates, while red arrows shows the direction of propagation of the light along with the plane of incidence. Adapted from Ref. 140. available components is presented in Figure 3.6.

Figure 3.6: Scheme of our MOKE apparatus set to measure polar Kerr rotations. Adapted from Ref. 142.

Components of the MOKE setup were mounted on a home-made breadboard made

of acrylic glass and a Al metal sheet. The light source consisted of a 532 nm YAG green

laser (Lotis TII). Light was polarized by a calcite Wollaston prism (Edmund, #68-

821) in either the p-ors-polarization before probing the sample. These prefixes come from the German senkrecht or perpendicular for s-polarized light, and parallelklang

81 or parallel for p-polarized light referring to the electric field E c component of the

incident plane polarized light.119 They could also be referred to as out-of-plane and

in-plane polarized light, respectively. All tested samples were deposited on NaCl

prisms and probed either as internal or external reflection. In order to effectively

concentrate the reflected light, it passed through a lens before reaching a second

polarizer (polarization filter 80 #G036325000 from Qioptic). Light was detected

using an Optic Ocean spectrometer (USB4000). The magnetic field was generated

with a C-frame electromagnet (GMW Associates, 45 mm #3470) and measured with

a LakeShore gaussmeter (475DSP model, 475-HMNT-4E04-VR) equipped with a Hall

probe (HMNT-4E04-VR).

The original goal of this experiment was to obtain Kerr rotation θK values for our

nanomaterials in order to quantitatively describe their magnetic behaviour. However,

the literature Kerr rotation values for the well-known 3d ferromagnetic elements Fe,

Co, and Ni are < 1◦ as show in Table 3.1. Measurements of such values would require an instrument with precise control over the polarizer’s angle of rotation – most likely an automatic controlled MOKE apparatus as opposed to our manual set-up. Therefore, an experimental alternative to measuring degrees of rotation was required in order to optically characterize our nanomaterials.

3.4. An alternative to direct measurement of the Kerr rotation

An alternative to measuring Kerr rotation θK values is to measure the intensity of the

reflected light with the sole purpose to study the relative change in intensity. This

is a valid alternative as the reflected light rotated any minute quantity as a result of

the presence of the Kerr effect could be blocked by strategically positioned polarizers

before and after sample reflection resulting in reflected-rotated light unable to reach

the detector.

82 Sample θK Specifics Reference (| deg |) Fe ∼0.32 Fe (fcc) on Au 143 0.82 Fe (bcc) at 1.61 T 144 Co 0.41 and 0.45 Co (hcp)and(fcc) at 2.8 T 145 0.70 Co (hcp)at1.61T 144 Ni ∼0.14 Ni (fcc) on Cu at 1.88 T 146 0.26 Ni (fcc) at 1.61 T 144

Table 3.1: Literature values for Kerr rotation of 3d iron-series ferromagnets.

For our purposes, instead of directly measuring the Kerr effect, it was decided that a practical methodology would be to record light intensity as the observable variable in our experiments. In order to achieve this, the polarizers (both the polarizer and the analyzer) had to be set within a small angle away from extinction (e.g., polarizations perpendicular to each other in a ∼0◦ to ∼90◦ arrangement allowing only a small

amount of light to pass through). A small angle from extinction was adopted as

an attempt to measure any small changes in the light intensity reflected from our

130,147 samples. The relationship between Kerr angle φK and light intensity is:

δ I − I0 φK = (3.3) 2 I0 δ is the fixed small angle at which the two polarizers are set to produce the minimum

light intensity of the reflected light, I is experimental light intensity measured, and I 0

130,147 is the average intensity equivalent to the intensity at φK =0 or no Kerr rotation.

3.5. Results and discussion

3.5.1. Preliminary optical experiments

Before conducting MOKE measurements, in order to test the optics of our home-

assembled setup, and our ability to prepare NPs, optical measurements were car-

ried out. Specifically, plasmon resonance studies were performed as a result of their

83 well-known effect on metal nanomaterials. Amongst our metal choices, Ag was our preferred alternative due to its mid-range cost between low-cost Cu and high-cost

Au. Additionally, Ag has a prominent optical activity and is readily utilized in signal amplification during surface-enhanced Raman spectroscopy technique.148–150

The two metal structures synthesized in our studies were gas-phase ligand-free NPs and thin films deposited onto substrates. Topographically these nanomaterials bound to a surface differ in that NPs are discrete species positioned some distance apart, while thin films are continuous structures. A detailed description of the instrument and justifications on our synthetic methods are found in the Experimental details chapter (i.e., Chapter 2). Briefly, within the nanogenerator, the formation of NPs oc- curs in the aggregation chamber. The aggregation zone houses a DC magnetron head that sputters the desired metal source generating an atomic vapour in the process.

These gas-phased atoms are then condensed in the presence of an inert gas cooling under vacuum conditions. Once formed, the NPs are swept out of the aggregation zone and collected onto a substrate at the deposition chamber. A modified version of this synthetic method was adopted to prepare thin films. By placing a collecting substrate directly onto the metal target at the center of the circular racetracks cre- ated by sputtering (i.e., away from the depletion area)151 material is deposited onto the desired substrate without further handling or fixing agents. For optical experi- ments, direct deposition of NPs and thin films onto NaCl prisms was implemented.

The benefits of using salt prisms are that NaCl is a transparent crystal77 allowing for transmission, while prisms enable to probe the metal sample directly at their exposed surface (i.e., external reflection) or through internal reflection with the light passing through the NaCl itself. Figure 3.7 shows the collection of Ag thin films.

A diagram of the optical set-up is shown in Figure 3.8. This set-up permitted probing of samples with either s-orp-polarized light. The interaction of light in either

84 Figure 3.7: (a) Photograph showing the sputtering head and the placement of a NaCl prism onto a Ag target. The prism must fit inside the circular groove produced by the plasma during metal etching. The prism holder consists of a paper clip completely wrapped in aluminum foil. (b and c) Diagram showing the side view of two different placement methods to rest the collection prism on the metal target. (b) At a 90◦ angle, the substrate collects an nonuniform density film on its hypotenuse face with its lower section collecting more material. (c) A hypotenuse-facing placement produced a more uniform density film. Method (c) was adopted. polarization will reveal the topographical composition of the sample. For instance, it is expected that a film consisting of discrete individual particles will behave differently when exposed to s-polarize light than to p-polarize light. This is due to the presence or absence of interparticle coupling.152

Interparticle coupling is affected by the plasma oscillations (i.e., induced dipole) of the metal NPs. Plasma oscillations (also called Langmuir waves) have a frequency

νplasma defined as:

 Ne2 νplasma = (3.4) Km0 where N is the density of free electrons, e is the electric charge per carrier, m is the mass of the charge carriers, K is the dielectric constant of the medium containing

153 the oscillating charges, and 0 is the permittivity constant. Thus, as the dielectric constant K of the surrounding increases, as would be the case for metal NP-metal

NP interactions, with s-polarized probing light compared to metal NP-air or metal

85 NP-substrate in the case of p-polarized probing light, the plasmon frequency νplasma

shifts to lower values.153 As a consequence of the relationship E = νh,whereh is

Plank’s constant, energy E also shifts to lower energies. In contrast, a film consisting

of an uninterrupted (continuous) material containing an abundance of electrons will

not differ in its interaction with polarized light behaving instead as a large single

entity.

Figure 3.8: Schematic of the optical set-up with a prism set for external reflection experiments. Light from a deuterium lamp (i.e., the light source) is polarized in either the s-orthep-polarization by the first polarizer. The light is then reflected by the sample and filtered through a second polarizer analyzer before reaching the detector. For visual aid, a prism positioned for internal reflection experiments is also shown on the upper right corner.

The NP sample preparation consisted of 3.7 nm Ag NPs deposited for 7 min on a NaCl prism. The conditions for the preparation of these NPs are presented on

Table 3.2, along with the mass spectrometer measurements in Figure 3.9. The NP arrangement was examined with both s-andp-polarized light as shown in Figure 3.10.

NP diameter Ar flow rate He flow rate IVDp (nm) (sccm) (sccm) (mA) (V) (cm) (Pa) 3.7 30.9 20 300 311.8 50 63.5

Table 3.2: Conditions for the generation of 3.7 nm Ag NPs. I is the sputtering current, V is the sputtering voltage, D is the distance of the metal target to the aggregation chamber exit, p is the pressure in the aggregation chamber during deposition.

86 Figure 3.9: Mass spectrum of 3.7 nm Ag NP diameter determined by in situ mass spectrometry. Line is Gaussian fit to the data.

There are a few observations obtained from Figure 3.10. First, deposition for 7

minutes forms a sub-monolayer film of NPs. The different spectra produced from the

two polarizations indicates that the film is made of discrete particles some distance

apart, and not a continuous film which could be obtained by longer deposition times

or by collecting materials directly on the target. Second, from the s-polarized spec- trum it is determined that s-polarized light produces a strong coupling effect among particles as the probing energy has shifted to lower energy values compared to the non-polarized light. This is the result of inducing a surface plasmon resonance phe- nomenon parallel to the surface and in the process coupling adjacents NPs. Third, from the p-polarized spectrum it is determined that the NP film resembles isolated particles (i.e., sub-monolayer film) suggestive by the higher energy shift. This is the result of inducing a surface plasmon resonance phenomenon above and below the sur- face resulting in uncoupled or non-interacting plasmon resonance. These interactions are graphically represented in Figure 3.11. The lower energy plasmon resonance peak

87 Figure 3.10: Absorbance spectra of 3.7 nm diameter Ag NPs deposited onto a NaCl prism. Red spectrum shows the interaction of the NPs with non-polarized white light. Green spectrum shows the interaction of the NPs with p-polarized light. Orange spectrum shows the interaction of the NPs with s-polarized light. at 2.37 eV is characteristic of surface in-plane dipole interactions among particles in the film, while the higher peak at ∼3.16 eV is characteristic of out-of-plane surface interactions.152,153

Experimentally it has been reported that dodecanethiol-capped Ag NPs of 3.5 nm

diameter dispersed in hexane solvent show a 0.25 eV shift to lower energy in their

absorbance peak when these NPs are close-packed via Langmuir-Blodgett films.153

Similarly, ligand-free Ag NPs of 3.5 ± 0.5 nm deposited onto glass showed a lower

peak shift of over 1.8 eV when the NP concentration on the substrate was increased

(i.e., by decreasing interparticle distance).152 Both examples utilized non-polarized

light to probe their samples while their peaks shifts were attributed to a physical

change to the films by reducing interparticle distance, which in turn increases par-

88 ticle coupling during absorption experiments. Our peak shift value of ∼0.8 eV was

achieved with light polarization only without further modifying interparticle distance.

A comparison of the peak shift with literature values, however, is not straightforward

without knowing the interparticle distances for our sample.

Figure 3.11: Diagrams illustrating the interaction of surface unmodified Ag NPs with polarized light. (a) The oscillating electric field E c of linearly polarized light inducing surface plasmon resonance on NPs. (b) s-polarized light couples particles with their nearby neighbours. (c) s-polarized light does not promote coupling among particles behaving instead as single particles.

Similar optical experiments were carried out with Ag thin films collected on-target

for 7 min on a NaCl prism. Its s-andp-polarized spectra are presented in Figure 3.12.

There are a few observations gathered from Figure 3.12. First, the material that self- deposits onto the salt prism at the metal target for 7 min forms a continuous film as seen by the similar s-andp-polarized spectra. Second, besides a higher intensity absorbance value for the non-polarized light spectrum, all three spectra show the same peak position. All three spectra confirm a peak position that does not shift

89 regardless of light polarization.

Figure 3.12: Absorbance spectra of an on-target Ag thin film deposited onto a NaCl prism. Blue spectrum shows the interaction of the NPs with non-polarized white light. Red spectrum shows the interaction of the NPs with p-polarized light. Green spectrum shows the interaction of the NPs with s-polarized light.

Although the topographical characterization of these samples could also be achieved using other surface characterization techniques, our intent was to test our optical equipment and design, which ultimately would allow us to perform MOKE experi- ments. For our purposes, this preliminary experiment validated our setup.

3.5.2. MOKE measurements

All MOKE measurements were performed in a polar configuration on materials de- posited onto NaCl prisms. Data were gathered in a dark room in order to avoid interference with ambient light. As previously stated, a MOKE setup in a polar con-

figuration results in a larger Kerr effect when compared to longitudinal or transverse

90 geometries,119,126 which also affects the intensity of the reflected light.

Initially, tests were performed under external reflection. However, as shown in

Figure 3.13, most of the reflected light was not reaching the detector as a result of diffuse reflection (also called scattering). For example, from visual inspection our

films consisted of a rough surface, most likely forming a polycrystalline structure based on the sputtering nature of the synthetic process. To correct for this issue, internal reflection was adopted at first.

Figure 3.13: Diagrams representing (a) specular reflection from a smooth surface where the reflected rays are all parallel to one another, and (b) diffuse reflection from a rough surface where the reflected rays travel in random direction. Adapted from Ref. 1.

3.5.2.1. Internal reflection

Experiments using internal reflection resulted in better signal acquisition as a result of a smoother surface at the salt-Ag interface. However, it was soon realized that during internal reflection the light has to pass through the salt prism twice: once before reaching the sample and once after reflection. Since our goal was to measure polarization rotation in the presence of a magnetic field, internal reflection in a salt prism added complexity to the experimental set-up. Although NaCl is diamagnetic, the presence of a magnetic field makes the material optically active which experi- mentally added a level of complexity to our signal as described in the Faraday Effect section. Due to the added signal complexity, experiments performed under the inter- nal reflection configuration were too difficult to interpret and thus not pursued any

91 further.

Additionally, although it was unbeknownst at the time, the source of magnetic

behaviour of our materials was the result of surface oxidation. Since the material at

the interior of the films were not exposed to air, it remained metallic Ag as proved in

later chapters of this thesis.

3.5.2.2. External reflection

In contrast to internal reflection, MOKE external reflection experiments only dealt

with a single material surface without potential magneto-optic effects from the sub-

strate. For these experiments, Ag thin films and NPs were exposed to s-andp- polarized light under a magnetic field H strength in the range of –0.25 to 0.25 T

(the upper and lower limits of our electromagnet in the absence of a cooling loop).

There were several modifications to the preliminary experimental setup as a result of problems encountered during data gathering. For instance, the issue of diffuse reflection was minimized by adding a converging lens to the reflected light in order to concentrate the signal just before reaching the detector. Additionally, for signal simplicity, it was decided that a single-wavelength light source in the form of a green laser (530 nm, 2.34 eV) produced a cleaner signal than the white light source used in the preliminary optical experiments. The MOKE experiments themselves consisted in monitoring the peak intensity positioned at 530 nm while exposing the sample to different magnetic field strengths. Laser light as a source for probing light is not un- common in MOKE experiments.130 The probing penetration depth occurs in the first

152 to 20147 nm of the sample, thus this technique is also known as surface MOKE or

SMOKE. As a result of the low surface penetration of the probing light it is expected

that the NaCl substrate would not play a role in the signal obtained from thin films

as they are continuous surfaces.

Furthermore, when incident linearly polarized light is reflected from any surface,

92 it becomes elliptically polarized. This phenomena, however, is eliminated if the in-

cident light is purely s-(i.e., perpendicular to the plane of incidence) or p-polarized

(i.e., parallel to the plane of incidence) retaining its initial polarization state upon reflection.147 If, however, the incident s-orp-polarized is reflected off a magnetized surface, the light becomes elliptically polarized. For example, if a p-polarized beam is reflected off a magnetic sample, the reflected light will have a major component in the p-plane with a minor component out-of-phase in the s-plane resulting in an overall elliptically polarized light (i.e., Kerr effect). In order to avoid this issue, a Wollaston prism is added to the MOKE set-up. A Wollaston prism is able to separate incoming light into purely s-andp-polarized light by approximately 1-cm apart on a horizontal axis allowing to test a single polarization at the time by adjusting the probing angle reaching the analyte. The Wollaston prism was placed between the sample and the polarizer closer to the light source.

MOKE measurements are presented in Figure 3.14. By monitoring the light inten- sity of the 530 nm peak and utilizing Equation 3.3, the Kerr angle φK was calculated

accordingly. As seen in this figure, the intensity of the monitoring peak was affected

by the different magnetic field strengths applied. However, the interpretation of these

values is not straight-forward.

Figure 3.14a describing the Ag thin films suggests that, based on the highest

signal change, the greatest Kerr angle φK is obtained at a magnetic strength of -

0.22 T during s-polarization resulting in a value ca. 0.3◦. Similarly, for Ag NPs,

◦ Figure 3.14b indicates the greatest φK of ca. 0.2 is obtained during s-polarization

at 0.1 T. Although this latter angle is a negative value in the figure, literature values

are presented as absolute values.143–146 Moreover, the optical preliminary experiments

had shown that NPs samples are composed of individual structures which suggests

that the signal has a contribution from the NaCl substrate. By comparison, the

93 Figure 3.14: Kerr effect measurements of Ag by monitoring the peak intensity at 530 nm of a green laser as a function of magnetic field strength. Measurements done in external reflection mode of Ag (a) thin film and (b) NPs deposited onto NaCl prisms.

literature value of the polar Kerr rotation for Ag in an applied magnetic field H of

1Tis∼0.004◦.154 The literature value is two orders of magnitude lower than our

experimental value indicating a large discrepancy. It is important to note, however,

that our samples are not pure Ag but, as proven latter on this thesis, a mixture of

Ag oxides and composed of nanodimensioned structures, and ultimately magnetic.

A concern for our choice of polar MOKE testing is that, in general, for magnetic

thin films there is an attempt to reduce their potential energy by shape anisotropy

which tends to align the magnetic moment μ of the film parallel to its surface. For

this reason, polar MOKE measurements require the presence of an external magnetic

field of sufficient strength in order to bring the magnetization M alignment perpen-

dicular to the surface of the film.126 Unfortunately, the maximum field strength in

our electromagnet reached ±0.22 T. In addition, as expected for noble metals, their magneto-optical activity is much weaker compared to ferromagnetic materials at mag- netic fields of ca. 0.5–1 T requiring in theory the application of magnetic fields in the several hundreds T.155

94 3.6. Conclusion

The light intensity change measured during MOKE experiments for both Ag thin

films and NPs provided a good indication that these samples were demonstrating magnetic behaviour. The positive response of these control experiments warranted a more systematic study of these nanomaterials with superconducting quantum in- terference device (SQUID) magnetometry. The fact that a positive MOKE response is linked to spin-orbit coupling and exchange interactions, further supported SQUID magnetometry studies.

There were several interesting and unclear aspects of these positive results which required further analysis, for instance, the presence of magnetism without surface assembled monolayers, which many literature sources named as a prerequisite to the induction of ferromagnetism in diamagnetic coinage metals. Additionally, the role of nanodimensionality needed to be addressed, likewise surface effects such as oxidation as a possible process for the magnetic behaviour. Ultimately, the question that needed to be answered was the origin of the observe unconventional magnetism found on these nanomaterials.

95 Chapter 4

Inducing ferromagnetic behaviour in copper

nanoparticles and thin films through

non-stoichiometric oxidation1

4.1. Introduction

Recent years have seen an increasing number of examples of ferromagnetic behaviour being induced in diamagnetic noble metals,26,73,156 along with an increasing amount of debate over the chemistry and physics responsible for the observed magnetism. The coinage metals (i.e. Cu, Ag, and Au) have a d10s1 electronic configuration and are well-known to be diamagnetic as a result of delocalization of the free electron density in the s band.2,30 Contrary to expectation, however, hysteretic magnetization loops have been reported for nanoparticles (NPs) of Au,38,40 Ag,40 and Cu40 as well as for nanometer thin films of Au157 whose surfaces were functionalized with self-assembled monolayers (SAMs) of sulfur-containing ligands.38,40,157 Crespo et al.38 described how

1.4-nm Au NPs behave ferromagnetically at room temperature and discussed the im-

1Adapted with permission from J. Phys. Chem. C 2006, 120, 7388-7396. Copyright 2016 American Chemical Society.

96 portance of capping with strongly binding dodecanethiol to induce magnetism. Sim- ilar results were observed by Suda et al.,78 who found that Au thin films modified with a SAM of 4-[6(hexyldithio)hexyloxy]azobenzene also exhibit ferromagnetic-like behavior. This latter study purposely demonstrated the importance that the SAM – more specifically, its molecular dipole moment – has on the observed surface magne- tization by monitoring a reversible modulation of the magnetism and work function while inducing a cis-trans photoisomerization of the azo moiety.

Although the origin of the observed magnetism in nanodimensioned noble metals is not well understood and is still being debated,26,73 many models invoke the presence of surface-bound molecules as conditional to the occurrence of magnetism.42,43,158,159

However, the requirement of capping molecules has been called into question by a set of recent experiments wherein ferromagnetism was observed in uncapped nanocrys- talline Au films prepared by cluster deposition.44

A second set of results that questions capping molecules as the source of magnetism in nanodimensioned materials comes from the study of oxide-metal interfacial sys- tems. For instance, Gamelin and co-workers have reported room-temperature ferro- magnetism in several diluted magnetic semiconductors.160–162 Using trioctylphoshine

2+ oxide-capped 0.48% Ni -doped SnO2, which is a paramagnet as freestanding nanocrys- tals, this group was able to induce ferromagnetic behavior by gently annealing under

163 aN2 atmosphere. Ferromagnetic deactivation could also be achieved when the an- nealing was performed under aerobic conditions.163 Activation of ferromagnetism was attributed to an increase in formation of grain-boundary defects upon nanocrystal- nanocrystal interaction, whereas deactivation occurred as a result of the passivation of oxygen vacancies when air annealing was performed.163 Closer to the system of interest here, Gao et al.45 showed that solution-synthesized air-annealed pure ∼26- nm CuO NPs exhibit room-temperature ferromagnetism, which was attributed to

97 the presence of oxygen vacancies, specifically at the NP surface. Further studies by

47 the same team demonstrated ferromagnetism of bulk CuO/Cu2Ocomposites and

48 CuO/Cu2O microspheres. In both of these cases, the ferromagnetic behaviour was attributed to the interface formed between the two oxides and explained in terms of

the indirect double-exchange model. 47,48 Indirect double-exchange describes the pro-

cess in which electron hopping –maintaining spin orientation in the process– between

two ions of different oxidation states, in this case Cu(I) and Cu(II), occurs via a

shared oxygen atom.

Additionally, these interfacial concepts also extend to include pure noble metal

films such as Au, Ag, Pd, and Pt films deposited onto strong magnetic substrates

such as Fe, Co, and Ni, which show ferromagnetic properties attributed to the nature

of the interface between the noble metal films and the underlying supporting metal

substrate.74 The postulated chemistry leading to magnetism in these interfacial ma- terials include: symmetry breaking at surface atoms; charge transfer by hybridization between the d-bands of the two films; and, spin-orbit coupling between the two d- bands.74

The notion of atypical electronic structures at interfaces leading to magnetism is

supported by theoretical studies that predict magnetic behaviour in small unmodified

Ag clusters. The most magnetic of these in the range of 2-22 atoms is predicted to be

99 the icosahedral-shaped Ag13 cluster. In this cluster, 4d orbital of the central atom overlaps with those of its 12 outer-shell neighboring atoms, resulting in a charge

99 transfer favouring the outer-shell atoms. Thus, the resulting Ag13 cluster contains

an inner atom with a magnetic moment of 0.35 μB and outer-shell atoms with a

magnetic moment of 0.39 μB/atom. Because 12 of the constituent Ag atoms are

surface atoms, they are, in effect, interfacial. Consistent with theory, experiments

164 on Pt13 clusters supported in zeolite were shown to display magnetic behaviour.

98 Notice that these results speak to magnetism in nonfunctionalized noble-metal NPs and underscore the relatively immature nature of the current state of understanding.

Although it seems clear that interfaces are important, the role of functionalization versus the role of size, noting that most magnetic structures observed have been nanodimensioned, is not yet well understood.

Ideally, a study to distinguish between size effects and interfaces on magnetism in coinage metals would control these two parameters independently. To this end, in the present work, Cu NPs and thin films were fabricated in the gas phase under vacuum conditions (i.e., free of capping ligands). With the ability to control the average size of sub-12-nm particles, the size dependence of magnetic properties in pristine Cu NPs was examined. Through oxidation of both nanomaterials, the introduction of oxide- metal interfaces in partially oxidized surfaces was observed to induce ferromagnetism, and the induction of ferromagnetism in nanomaterials of pure metallic Cu could be correlated with nonstoichiometric oxidation at the surface. The systems are complex, however, with further oxidation resulting in a decrease in their ferromagnetic signal.

4.2. Experimental

The specific experimental methodology for Cu nanomaterials including synthesis and collection are found on Section 2.10.

4.3. Results

4.3.1. Cu NPs

The size distribution of NPs produced in the gas phase were determined by mass spectrometry and are illustrated in Figure 4.1a. Source conditions were adjusted to obtain NP size distributions with mean diameters of 4.5 ± 1.0 nm, 6.5 ± 1.0 nm, and 9.0 ± 1.8 nm. Efforts to obtain smaller nanoparticles to increase the size range

99 studied were unsuccessful. We speculate that by setting the mass spectrometer to

filter a specific NP diameter, rather than (as it was used here) to monitor synthesized

species, both size and monodispersion could be improved considerably. Additionally,

it is possible that the operational lower limit of the instrument was reached with our

applied parameters. These three source conditions presented in Table 2.2 were used

to generate and deposit all of the NPs studied in this work. For each of these source

conditions, NPs were deposited for 30 min onto NaCl substrates and stored in sealed

containers under Ar until ready for SQUID analysis. Note that a linear diamagnetic

response has been subtracted from all hysteresis loops. Figure 4.1b shows an example

of a diamagnetic correction performed on raw SQUID data. SQUID magnetometry

data of these NPs are presented in Figure 4.1c. As can be seen, the M vs μ0H loops indicate a clear magnetic response.

The inset in Figure 4.1c is a closer inspection of the loop near the origin displaying

a weak coercivity and remnant magnetization (e.g., μ0Hc = 25.7 mT and Mr/M1T = 9.20 %, respectively for 4.5 ± 1.0 nm particles), which is a clear indication that the Cu NPs behave as a soft ferromagnet at room temperature. No significant size effect can be observed over the investigated size range, however, and all three sizes of NPs show similar hysteresis loop (i.e., shapes) and saturation magnetization (i.e., intensities).

To better understand the impact of oxidation on magnetic properties, a number of NP samples (mean diameters of 4.5 ± 1.0 nm) were deposited onto NaCl for a total of 30 min and stored under either Ar or O2 atmospheres for 6 days. The data in Figure 4.2 are SQUID data collected for the two types of samples. As can be seen,

exposure to oxygen caused a significant decrease in the saturation magnetization. The

lower inset shows a close up of the origin; both coercivity and reduced remanence are

virtually unchanged between the samples stored in Ar or O2 gases, indicating a loss

100 Figure 4.1: a Cu NP diameters determined by in situ mass spectrometry. The re- sponses have been shifted vertically and the response for the smaller NPs has been multiplied by a factor of 1.5 for ease of visualization. Diameters are 4.5 ± 1.0 nm (), 6.5 ± 1.0 nm (×), and 9.0 ± 1.8 nm (). Lines are Gaussian fits to the data. b Sample diamagnetic correction performed on raw data from a Cu film kept in Ar. Solid trend represents the raw data and the dashed trend is the calculated linear diamagnetic correction. c SQUID measurements of Cu NPs. Inset is close up of the origin. Lines are guides to the eye.

101 of magnetic centers, not a change in type of magnetic centers. The upper inset shows

an AFM image of Cu NPs, illustrating the monodispersed topography of the sample.

Interestingly, the NP size based on AFM is ∼10 nm. This could result from the well- known lateral resolution limitations of AFM due to the cantilever tip size and sample roughness when trying to distinguish between two distinct NPs165,166 (in our case sub-

10 nm particles). This is compounded by the high NP density of our samples, which

could lead to spontaneous restructuring, for example, due to capillary condensation

between particles,166 in an attempt to reduce surface energy among particles.167

Figure 4.2: Room temperature SQUID magnetometry measurements of Cu NPs stored in either Ar ()orO2 () atmospheres. Bottom inset shows a close up of the origin. Lines are a guide to the eye. Top inset shows an AFM image with a 50 nm bar.

4.3.2. Cu films

At first glance, the results obtained suggest that pristine NPs are ferromagnetic, that there is no size dependence to the magnetism, and that prolonged exposure to oxygen attenuates the ferromagnetic behaviour. To investigate further whether the

102 ferromagnetic properties were broadly size-specific to sub-12 nm particles, the mag- netic properties of Cu films prepared under identical source conditions, but collected on Si substrates at the target (location b in Figure 2.12) for a total of 10 min, were also examined. SQUID magnetometry and AFM data are shown in Figure 4.3. From

AFM, it can be seen that the Cu films had a very coarse topography and consisted of relatively large faceted crystals with dimensions exceeding 100 nm. From the SQUID data, it can be seen that Cu films displayed the same ferromagnetic behaviour as the Cu NPs. Accordingly, we observe Cu films behaving ferromagnetically, which is puzzling because bulk Cu is diamagnetic. Also, we observed no evidence of size- dependent ferromagnetism because both NPs and thin films behaved similarly. As was observed for the Cu NPs (Figure 4.2), exposure to O2 diminished the saturation magnetization of the thin-film samples.

To further explore the effect of oxidation, several Cu films deposited on Si were heat-treated in air (i.e.,50◦C for 30 min, 70 ◦C for 90 min, and 100 ◦C for 11 hr).

The temperatures and times were selected such that the magnetization was roughly halved between each subsequent treatment.

The SQUID traces shown in Figure 4.4 illustrate how the saturation magnetization steadily decreased with increasing temperature. The results are analogous to those seen in Figures 4.2 and 4.3 for Cu NPs and films, respectively, which collectively indicate that oxidation leads to a decrease in ferromagnetism. Accordingly, this trend would suggest that the pure Cu NPs, as first produced in the vacuum apparatus and prior to any oxygen exposure, would have an even higher saturation magnetization.

4.3.3. Cu NPs - in situ MQCM measurements

To measure the magnetic properties of pristine NPs under vacuum, Cu NPs were deposited onto a QCM substrate, and their qualitative magnetic behavior (i.e., mag- netic vs nonmagnetic) was evaluated using the MQCM technique. Before describing

103 Figure 4.3: SQUID magnetometry measurements of Cu films stored in either Ar () or O2 () atmospheres. Bottom inset shows a close up of the origin. Lines are a guide to the eye. Top inset shows an AFM image with a 500 nm bar. our MQCM results, however, the proximity effect of the bar magnet to the QCM crys- tal when performing this technique are presented in Figure 4.5a during an MQCM measurement of a black QCM crystal in which a bar magnet exerts a 0.1 T magnetic

field to it. Accordingly, when the bar magnet was brought into close proximity of the the blank QCM to exert a field, a decrease in QCM oscillation frequency was observed, which would indicate a non-magnetic behaviour as described next.

The general MQCM trends for different magnetic materials can be found in Fig- ure 4.5b. This figure illustrates the effect of ferromagnetic and weakly paramagnetic materials in filing form on the face of the QCM. As can be seen, weakly paramagnetic

Al has little effect and both Al and the blank QCM display an analogous decrease in oscillation frequency with increasing magnetic field strength. Ferromagnetic Fe, on the other hand, produces a positive frequency change. The notable trend in Fig- ure 4.5b is that of ferromagnetic materials which resulted in a frequency increase during MQCM experiments (i.e. positive Δf MQCM). MQCM measurements are in-

104 Figure 4.4: SQUID magnetometry measurements of heat-treated Cu films. Films were heat-treated in air at 50 ◦C(), 70 ◦C() and 100 ◦C(). An as-prepared sample stored at room temperature under O2 is also shown ().

herently convoluted because current flow through the QCM crystal is itself sensitive

to the presence of magnetic fields and to proximity effects. 110,168

Two plausible explanations for the positive Δf MQCM observed when probing a ferromagnetic material during MQCM measurements can be found in the original

Janata et al. paper.109 The “acoustic impedance perturbation” hypothesis is based on increased magnetic coupling between magnetic particles in a film in the presence of an applied magnetic field H. This magnetic coupling increases the stiffness of the

film and, together with the coupled motion of particles during a regular QCM ex- periment, promotes the propagation of the acoustic waves resulting in a decreased acoustic impedance and in an increased frequency.109 This hypothesis however does

not explain the experimental negative Δf MQCM we observed with Al. Although, in our setup a proximity effect (i.e., proximity of the bar magnet to the oscillating QCM)

could explain this behaviour as a negative Δf is observed with a blank QCM. The

second hypothesis borrows from the principle of vibrating sample magnetometry fun-

105 Figure 4.5: (a) Proximity effect of a bar magnet during MQCM raw measurements of a blank QCM crystal. The bar magnet is positioned to exert a 0.1 T magnetic field onto the QCM crystal. (b) MQCM measurements of ferromagnetic and weakly-para- magnetic materials. Blank QCM (), Al fillings () and Fe fillings ()exposedto ≤0.1 T magnetic fields with a bar magnet in an MQCM experiment. Lines are guides to the eye.

damental in SQUID magnetometry. The “magnetic disturbance” hypothesis describes

how the particles in motion create a magnetic field gradient which disturbs the uni-

formity of the applied magnetic flux density.109 This field disturbance is transduced

as an oscillation and picked up by the quartz crystal, in turn interfering with its own

oscillation in a constructive or destructive fashion resulting in a positive or negative

109 Δf MQCM respectively. The latter magnetic disturbance hypothesis is able to ex-

plain our experimental positive and negative Δf MQCM and is in a better qualitative accordance with our own experimental observations.

During in situ MQCM measurements of nascent Cu NPs under vacuum conditions, negative frequency changes were observed when the sample was exposed to a magnetic

field (Figure 4.6a). Furthermore, the frequency was comparable to that observed for the blank QCM prior to deposition of Cu NPs on the QCM face (Figure 4.5). The lack of effect in the presence of Cu NPs is indicative of the pristine NPs being a

106 diamagnetic material, as would be expected for bulk Cu. Note that these data also

suggests parasitic magnetic impurities are not responsible for the observed magnetic

response, as a magnetic response would have been observed in this case. As can be

seen in Figure 4.6, however, upon the introduction of air into the vacuum system,

there was a positive response to the magnetic field that manifested itself as a decrease

in the magnitude of the frequency change – albeit still negative. Oxidation was

further induced with several ex situ heat treatments, as the magnitude of the negative frequency change was steadily reduced, suggesting that uptake of oxygen by the NPs induces magnetism. In contrast to the data above, the MQCM results demonstrate an increase in magnetic behaviour as a result of oxidation of the Cu NPs.

Figure 4.6: Frequency change measurements during an MQCM experiment of Cu NPs with increasing mass of oxygen. Down arrow (up arrow) represents the bar magnet at a distance exerting a ∼0.1 T (negligible) magnetic field. Panel a represents a as-prepared sample under vacuum, panel b shows a gain of 0.15 μg, while panel c represents a gain of 0.95 μg. Ticks on the abscissa represent 100 s.

An advantage of the MQCM technique is that the QCM crystal is a microbal- ance that permits the simultaneous measurement of the magnetic properties and the quantification of the mass of oxygen taken by the Cu NPs. In Figure 4.7a, a plot of the frequency-change response of the Cu NP-coated QCM against the mass of oxy-

107 Figure 4.7: a. MQCM experiment of nascent Cu NPs () with subsequent oxidation with ambient air () and heat-treatment (). All measurements performed at room temperature. Line is a guide to the eye. b. Saturation magnetization from SQUID measurements of heat-treated Cu films. Line is a guide to the eye. Inset shows SQUID measurements from 0 to 1 T of the Cu films.

gen uptake (as determined from the absolute QCM frequency) by the NPs is shown.

Initial exposure to oxygen caused a positive change in the oscillation frequency (i.e.

Δf MQCM became less negative). The trend completely reversed, however, near 1 μg of oxygen taken up for 62 μg of Cu. After this point, further absorption of oxygen manifested itself as a steady decrease (increasingly negative Δf MQCM)intheQCM oscillation frequency response to the presence of a magnetic field. That is, with oxy-

gen uptake, the sample initially became magnetic, but as oxygen pickup continued

past a certain threshold, the Cu NPs became less and less magnetic. To illustrate the

trend more clearly, the SQUID saturation magnetization values of the heat-treated

Cu films shown in Figure 4.4 are presented alongside the MQCM data in Figure 4.7b.

In this case, as the films were exposed to higher temperatures, additional oxidation

took place, leading to a further reduced saturation magnetization.

To test for potential impurities and to analyze the oxidation state of the Cu films,

XPSanalysiswasperformedonseveralsamples.TheCufilmskeptinArandO2,

108 as well as the heat-treated samples shown in Figures 4.3 and 4.4, were all analyzed by XPS. Sample surveys of Cu films kept in O2 and Ar atmospheres before and after in situ sputtering are shown in Figure 4.8. Surveys show samples of highly pure Cu with no measurable amounts of ferromagnetic elements (i.e., Fe, Co, or Ni). Other elements such as C, O, and S were also present. Interestingly, S was absent in all heat-treated samples and sputtered surfaces, suggesting that it was only physisorbed on the O2- and Ar-kept samples.

Figure 4.8: Peak assignments for the complete XPS surveys of O2- and Ar-kept Cu film before and after sputtering.

109 Figure 4.9 shows the S region of the O2-keptCuthinfilmsamplebeforeandafter sputtering, as well as the heated sample at 50 ◦C.

Figure 4.9: XPS data showing the S region of the O2-keptCufilmbeforeandafter sputtering, as well as the Cu film heat-treated at 50 ◦C.

Accordingly, no evidence of an impurity effect was observed, and the ferromagnetic signature observed at room temperature can be attributed to the Cu films, specifically.

◦ Furthermore, the films kept in Ar and O2, and the sample heated at 50 Cwere sputtered in situ inside the XPS instrument for analysis. Fresh surface was exposed after removing of 50 nm of material. To study the oxidation state of Cu, the Cu 2p3/2 peak was monitored.

XPS data showing the Cu 2p region of Cu films kept in O2 and Ar atmospheres before and after in situ sputtering are presented in Figure 4.10a with their corre- sponding O 1s region in Figure 4.10b.

Similarly, XPS data showing the Cu 2p region of Cu films heat-treated at 50 ◦C,

70 ◦C and 100 ◦C are presented in Figure 4.11a with their corresponding O 1s region in Figure 4.11b.

110 Figure 4.10: XPS data showing the (a.)Cu2p and the (b.)O1s regions of the O2- and Ar-kept Cu film before and after sputtering. Raw data are shown as scattered plots. All amplitudes are maximized for ease of visualization, thus are not comparable. a. Peak fitting and Shirley background are shown for the Cu 2p3/2 peak. b. Peak fitting (solid and dotted lines) and Shirley background are shown for the O 1s region. C-OandC=OpeakareasintheO1s region match those corresponding to the C 1s region (latter data not shown).

111 Figure 4.11: XPS data showing the (a.)Cu2p and the (b.)O1s regions of Cu films heat-treated at 50 ◦C, 70 ◦C and 100 ◦C. Raw data are shown as scattered plots. All amplitudes are maximized for ease of visualization, thus are not comparable. a. Peak fitting and Shirley background are shown for the Cu 2p3/2 peak. The appearance of the shakeup satellite is indicative of Cu(II) presence. b. Peak fitting (solid and dotted lines) and Shirley background are shown for the O 1s region. C-O and C=O peak areas in the O 1s region match those corresponding to the C 1s region (latter data not shown).

112 Because the binding energies of the Cu 2p3/2 peak for both Cu(I) and Cu(0) over- lap within the standard deviation of each other,169 the modified Auger parameter was used for characterization.169,170 An Auger electron is created when a core-level hole created after the photoejection of a core electron (as in the case of XPS analy- sis) is filled with an outer-shell electron while simultaneously ejecting a third Auger electron. The modified Auger parameter (α’) is defined as the sum of the kinetic energy of the Auger electron (EK ) and the binding energy of the photoelectron (EB),

170–172 α’=EK +EB. It is of great aid in determining the chemical state of materi- als;170 accordingly, the modified Auger parameter was utilized to identify the present

Cu species more reliably. Our results and literature comparison are provided in Ta- ble 4.1. Based on literature average peak positions,169 all non-sputtered films were determined to be composed of surface copper oxides with metallic Cu underneath after sputtering, with the exception of the 50 ◦C heat-treated sample. This latter result is not surprising, as the oxidation layer is expected to thicken (due to oxygen diffusion) with temperature.173

To quantify the different Cu species present in our films, the method described by

Biesinger et.al. was utilized.169,174 This method does not make a distinction quan- tifying Cu(0) and Cu(I) as separate species; thus, their aggregated percentage is presented as a ratio in Figure 4.12. In the same figure, the atomic percentage of

Cu(II), quantified with the help of the shake-up satellite peaks in the Cu 2p region

(Figure 4.11a), and of O are also shown. The curve fitting of the O 1s region for all samples is shown in Figures 4.10b and 4.11b.

113 Figure 4.12: Atomic percentages of Cu and O via XPS in several Cu films. Cu(0) and Cu(I) species (), Cu(II) (), and O (). Dark shaded areas represent sputtered samples while light shaded areas are heat-treated samples.

Sample Cu 2p3/2 Modified Auger Reference (eV) parameter (eV) Cu(0) literature avg. 932.61 ± 0.21 1851.23 ± 0.16 169 Cu(I) literature avg. 932.43 ± 0.24 1849.19 ± 0.32 169 Cu(0) experimental 932.63 ± 0.025 1851.24 ± 0.025 169 Cu(I) experimental 932.18 ± 0.12 1849.17 ± 0.03 169 Ar-kept 932.74 1849.20 this work Ar-kept + sputtering 932.70 1851.10 this work O2-kept 932.58 1849.57 this work O2-kept + sputtering 932.76 1851.10 this work 50 ◦C treatment 932.76 1849.42 this work 50 ◦C treatment + sputtering 932.69 1850.52 this work 70 ◦C treatment 932.47 1849.28 this work 100 ◦C treatment 932.51 1849.28 this work

Table 4.1: Literature comparison to our results of the Cu 2p3/2 peak positioning and the modified Auger parameter.

114 4.4. Discussion

4.4.1. Size dependence

Size-dependent properties are a well-known characteristic of NPs and are highly de-

bated as contributors to the unconventional magnetism that has been observed in

nanodimensioned coinage-metal systems. The sizes of NPs studied in this work were

well-defined by mass spectrometry, as can be seen in Figure 4.1a, and ranged from

as little as 3 nm to as large as 12 nm with mean diameters of 4.5 ± 1.0 nm, 6.5 ±

1.0 nm, and 9.0 ± 1.8 nm. For all of these size ranges, SQUID magnetometry data

(Figure 4.1b) were very similar for all samples, displaying coercivity and remanence

values that are characteristic of soft ferromagnetic materials. Thus, there is no evi-

dence of size correlating to observed magnetic properties over the investigated range.

We note, however, that thiol-capped Cu NPs with a 2.3 nm diameter were found to ex-

hibit a significantly stronger magnetization. 40 It is not clear whether this discrepancy

arises from the presence of capping molecules or the further reduced dimensionality.

Similarly, a comparison of the SQUID data in Figures 4.2 and 4.3, for Cu NPs and

Cu films, respectively, reveals very similar saturation magnetizations, coercivities,

and remanences, indicating that ferromagnetism can be observed in nanostructures

with sizes as large as ∼500 nm (see AFM inset in Figure 4.3). MQCM experiments, in which changes in magnetism were monitored in situ as source conditions were varied, revealed no indication that certain diameters of NPs were more “magnetic” than others. In fact, as noted above, pristine NPs generated under vacuum condi- tions displayed no magnetism at all, regardless of size (Figure 4.7a). As discussed below, the ferromagnetic properties observed cannot be specifically associated with nanodimensionality of the Cu metal component of these systems.

115 4.4.2. Oxidation

MQCM data of pristine NPs generated in a vacuum and then exposed to air, shown in

Figure 4.7a, demonstrate a rapid increase in magnetic response during initial oxidation

of the NPs but then a decrease with further oxidation by heat treatment in air. The

trend follows through in the SQUID data (Figure 4.7b), with a decline in saturation

magnetization upon heat treatment in air. By contrast, XPS analysis of the oxidation

states of several Cu films shown in Figure 4.12 revealed a decline in low oxidation

states of Cu [i.e., Cu(0) and Cu(I)] and a steady increase in oxygen content with

increased heat-treatment temperatures. Also, Cu(II) is a minor contributor to the

overall makeup of heat-treated films. This oxidation trend is consistent with previous

173 studies for the formation of Cu@CuOx core-shell NPs.

The immediate onset of magnetism upon exposure of pristine Cu NPs to air,

as can be seen in Figure 4.7a, is most consistent with a surface oxidation process

being responsible for generating magnetic centers. The fact that sputtering of Cu

films effectively removes all oxide and generates ∼90% pure Cu surfaces, as can be seen in the XPS data shown in Figure 4.12, is completely consistent with surface oxidation being the dominant process with interstitial oxidation within the interiors of nanomaterials being negligible.

4.4.3. Our magnetic system compared to other systems

In general, according to our data (Figure 4.12), there is no obvious evolution of a specific Cu oxidation state consistent with the rising and falling nature of ferromag- netism displayed in Figure 4.7a. However, others have found certain ratios of specific oxides in CuO/Cu2O microspheres, showing a maximum saturation magnetization

48 when the ratio reaches 73% CuO and 27% Cu2O. In contrast, our maximum satu- ration magnetization is found in Cu NPs and films (Figures 4.2 and 4.3, respectively)

116 kept under Ar or O2 with an undetectable CuO presence according to our XPS analy- sis (Figure 4.12). The absence of Cu(II) in our unheated samples is not surprising, as

it was previously found that surface formation of a layer of native Cu2O protects the

175 rest of Cu material against further oxidation. Moreover, conversion of Cu to Cu2O is more facile than conversion to CuO.176 This is due to O diffusion into tetrahedral

sites in fcc Cu leading to lattice expansion in the case of cubic Cu2O, as opposed to atom rearrangement during the formation of monoclinic CuO..176

This is illustrated in Figure 4.13 which shows the cell structures of Cu and its

common oxides Cu2OandCuO.

Figure 4.13: Unit cells of metallic fcc Cu and its common oxides cubic Cu2Oand monoclinic CuO. CuO adapted from Ref 177.

Another example in which oxidation plays a role in magnetization consists of 10-

100 nm nanocrystalline powders of Cu2O1+x where x>0 has been observed to show hysteresis connected to the appearance of Cu+2 having a 3d9 electronic configura-

tion;79 however, the authors stated that the reason for the ferromagnetism is unclear.

Interestingly, computational studies have proposed to explain ferromagnetism in

copper oxides with specific atomic ratios.49,50 In our case, Cu/O atomic ratios were

calculated from experimental MQCM and XPS data.

Assuming all NPs were 4.5 nm in diameter (based on average experimental results

of smaller NPs) and were perfect spheres, it can be estimated that each NP contained

NP × 3 Nttl =5.43 10 Cu atoms (based on an atomic radius rCu of 128 pm).

117 The volume of the last atomic shell VOS on a NP of radius rNP can be estimated as

4πr3 4π V = NP − (r − 2r )3 (4.1) OS 3 3 NP Cu NP × 3 from which it can be estimated that there are Nsurface =1.65 10 surface Cu atoms /NP.

Cu The total number of surface Cu atoms Nsurface in a 62.022 μg Cu sample (containing

Cu Nttl )isthengivenas

NP Cu Nsurface · Cu × 17 Nsurface = NP Nttl =1.79 10 (4.2) Nttl Knowing this sample uptakes 0.9471 μg of oxygen from QCM measurements (i.e.

O × 16 Nttl =3.56 10 ) and assuming these O-atoms will react at the surface, the ratio of surface Cu atoms to O atoms is then

N Cu 1.79 × 1017 surface = =5 (4.3) O × 16 Nttl 3.56 10 Thus, the resulting Cu:O atomic ratio was determined as 5:1.

Similarly from XPS analysis (Figure 4.12), we estimated a Cu/O ratio of approxi-

mately 3:1 (i.e., 3.2:1 and 2.9:1 for the Ar- and O2-kept Cu films, respectively). Com- paring our values with theoretical studies, ferromagnetism is believed to be induced

in a 48-atom Cu2O supercell model by interstitial oxygen defects at both octahedral (i.e., an interstitial O atom surrounded by nearest six Cu atoms or a 6:1 Cu/O ra- tio) and tetrahedral (i.e., a 4:1 Cu/O ratio) sites.49 In contrast, theoretical studies

with planar Cu/O clusters found that 4:1 and 4:2 behave diamagnetically, whereas

4:5 cluster (together with nonplanar 16:15 and 28:27 clusters) are ferromagnetic.50

Although both theoretical studies cite spin density localization and polarization as

the source of ferromagnetism,49,50 and although the compositions are not in accor-

118 dance with each other, they suggest that specific nonstoichiometry plays a role in the

magnetic behaviour.

Accordingly, the overall description of our ferromagnetic Cu system does not show

a size dependency, which is evident by the mass spectrometric sizing of NPs and their

corresponding SQUID measurements. Furthermore, the pure Cu NPs (under vacuum)

display no magnetism according to the in situ MQCM experiments. The ferromag-

netic behaviour in NPs appears only after exposure to oxygen. Cu thin films prepared

in a similar fashion corroborate the lack of size dependency and the ferromagnetic

behavior after oxygen exposure with further oxidation decreasing the saturation mag-

netization in both nanomaterials. Based on MQCM calculations with Cu NPs, the

highest magnetic signal was obtained when a 5:1 Cu/O atomic ratio was achieved.

Similarly, XPS analysis of Cu thin films showed higher magnetization when the ratio

was ∼3:1 Cu/O. XPS also provides evidence of surface oxidation while maintaining a metallic Cu interior. All this evidence reveals a clear trend in magnetism for Cu nanomaterials as a function of oxidation in which partial oxidation is key to maximiz- ing the ferromagnetic signal. Based on our MQCM and XPS evidence, we propose an intermediate step during the oxidation reaction of fcc metallic Cu to cubic stoi-

77 chiometric Cu2O (both of which are non-magnetic ) in which intercalating oxygen atoms form a nonstoichiometric oxide resulting in the observed magnetic response

(Figure 4.14). Although previous studies on the early stages of Cu oxidation involve

nucleation and growth of metal oxide nanoislands178–180 (as opposed to a uniform

oxide distribution), these studies were performed at 350 ◦C and higher temperatures.

On the contrary, our samples showing the highest magnetization were prepared at

room temperature. It is known that oxidation temperatures will affect O diffusion,

interfacial strain,and surface and interface energies, all of which play a role in the

nucleation and development of oxide morphologies.178 Thus, at low oxidation tem-

119 peratures, diffusion and strain, for example, should be at a minimum, resulting in surface oxidation and a cubic lattice (albeit Cu bonds are expected to elongate when

O is incorporated into the unit cell).

non-stoichiometric Cu (fcc) Cu O (cubic) Cu O 2 non-magnetic m n non-magnetic magnetic

Figure 4.14: Proposed oxidation stages of Cu from in its native fcc cell to cubic Cu2O. Both metallic Cu and Cu2O are non-magnetic.

We also propose a mechanism during the early stages of oxidation of Cu nanoma- terials in which magnetic centers are formed during partial oxidation of the material with decreasing magnetic response upon further oxidation (Figure 4.15). Because the in-vacuum Cu NPs behaved non-magnetically during the in situ MQCM experiments, it is also expected that the thin films will behave identically. This is collaborated by the fact that both nanomaterials show the highest magnetization with partial oxida- tion and declining magnetic signal upon further oxidation.

4.5. Conclusion

In conclusion, we have demonstrated the fabrication of highly pure Cu NPs and thin

films that behave ferromagnetically at room temperature. Our findings add valuable information to the current debate on the requirements to induce magnetic behaviour into (diamagnetic) coinage metals. The gas-phase procedure for the preparation of our nanomaterials enables the separation of the effects of size and surface chemistry on the possible impact in inducing magnetic behaviour. We have shown that there is no size effect on the observed magnetism; however, our nanodimensioned systems

120 Cu

Cu2O magnetic center

partial oxidation

nanoparticles in vacuum after air bled in chamber further oxidation OR in air OR heat treatment thin films

decreased magnetic non-magnetic magnetic response

Figure 4.15: Proposed growth mechanism for the early stages of Cu oxide on nanopar- ticles and thin films, and their observed magnetic behavior. are dominated by surface chemistry and surface effects. Moreover, it is clear that magnetism in our systems comes from non-stoichiometric oxidation. We propose this process as a fundamental source of magnetism in all Cu nanodimensioned systems and potentially all noble- metal systems.

121 Chapter 5

On the origin of the ferromagnetic signature in

silver nanoparticles and thin films1

5.1. Introduction

It has been recently observed that nanoscaled materials displayed new physical prop- erties not observed in their bulk counterparts; a striking example is the observa- tion of ferromagnetic-like behaviour in nanomaterials that are diamagnetic (i.e.,non- magnetic) in the bulk.26,73,156 The elements subject to such studies are, for the most part, the coinage metals (i.e., Cu, Ag and Au).40,181,182 The electronic valence con-

figuration of all coinage metals is d 10 s1 resulting in an overall magnetic moment in unbound individual atoms – the result of the unpaired s electron. However, as atoms form bulk materials, delocalization of the electronic states results in their well-known conductive and diamagnetic characteristics.77,183,184

Ferromagnetism in nanoscale coinage and noble metals was achieved through chemical surface modifications (i.e., capping ligands), as well as in uncapped (and undoped) nanomaterials and clusters. For instance, hysteretic magnetization loops

1This work has been submitted to Physical Chemistry Chemical Physics.

122 have been reported for coinage and noble-metal nanoparticles (NPs) of 2.3 ± 0.3 nm

Cu;40 2.3 ± 0.3 nm Ag;40 and, 1.4 nm38 and 1.9 ± 0.2 nm40 Au modified at the sur- face with self-assembled monolayers (SAMs) of sulfur-containing ligands. Non-sulfur containing SAMs of poly-N-vinyl-2-pyrrolidone have been shown to induced ferromag- netism in ∼3 nm Au and Pd NPs.37 Au thin films also show ferromagnetic behaviour when modified at the surface with alkanethiols (e.g.,C2-andC8-alkanethiols, and

157 78 C22-polyalanine) and azobenzene disulfide SAMs. Conversely, nanomaterials with unmodified surfaces such as Au thin films have been found to display paramagnetism106 and ferromagnetism.44 In the paramagnetic case, 27 nm thick Au films deposited onto cylindrical Pyrex substrates (i.e.,ring- shaped films), exhibited strong magnetism at 5 K but not at 300 K.106 The different magnetic responses were attributed to the Wigner radius r s (defined as the radius of the sphere occupied by an electron) which in a thin film would result in a large dipole layer when positioned at the surface aided by an almost-featureless textured

Au films.106 In the ferromagnetic case, the Au film was prepared by deposition of

∼2.6 ± 0.6 nm diameter clusters.44 The source of ferromagnetism is attributed to the existence of a core–shell nanostructure constituted of a bulk–like diamagnetic core and a magnetically active shell caused by the reduced coordination of surface atoms helped by a strong spin–orbit coupling in Au.44

Furthermore, predicted magnetic states have been theorized for nanoclusters of

Ag75,99 and Au75 – interestingly, these nanoclusters studies found the icosahedral

13-atom cluster (diameter of 0.548 nm for Ag185 and of 0.7 nm for Au186)tobe

75,99 the most magnetic with a moment of ∼5 μB per cluster. In the Ag case, the 4d orbital of the central atom in Ag13 overlaps with those of its 12 outer-shell neighboring atoms resulting in a charge transfer favouring the outer-shell atoms.99 Calculations for highly symmetrical Ag and Au clusters up to a maximum of 147 atoms (equivalent

123 to 1.664 nm radius for Ag185) has been reported with several degrees of magnetism.75

Since both Ag and Au atoms are d 10 s1 metals, it is claimed that the origin of this

magnetism in these larger clusters is due to spin alignment of the s-electrons from atoms in the outer shell in a “superatom”. 75 Additionally, experimental measurements

for icosahedral Pt13 (diameter of 0.7 ± 0.1 nm) supported in zeolite resulted in a magnetic moment of 3.7 ± 0.4 μB per cluster with only 15-20% of Pt atoms in this configuration contributing to the overall magnetic signal. 164

However, extrinsic sources of magnetism (e.g., impurities,187 or surface species

differing from the nominal composition of the material94) should not be dismissed.

Surface oxidation has been found to be another process inducing ferromagnetism in

otherwise diamagnetic materials when the formation of non-stoichiometric structures

is favoured such as in mixed oxides where interfaces are prevalent or in nanomaterials

where the percentage of surface atoms is high. For instance, Cu oxides in the form

45 47 48 of ∼26 nm CuO NPs, Cu/Cu2O composites, and Cu/Cu2O microspheres can behave ferromagnetically. Nanocrystalline powders of 10-100 nm size composed of

79 Cu2O1+x where x>0, also behave ferromagnetically. Theoretical studies of defect

formation in a 48-atom Cu2O supercell have shown that interstitial O surrounded by six (i.e., octahedral site) and four (i.e., tetrahedral site) Cu atoms49 induce ferro-

magnetic behaviour. Although this latter study was performed on a stoichiometric

Cu2O supercell, their proposed Cu/O atomic ratios resulting in ferromagnetism are nonstoichiometric in nature (i.e., 6:1 Cu/O for octahedral and 4:1 for tetrahedral).

Experimentally, we have recently shown that Cu NPs and thin films behave ferro- magnetically through nonstoichiometric oxidation of the surface with a calculated atomic ratio of ∼3-5:1 Cu/O.94 In our Cu study we conclusively showed that slight

surface oxidation leading to an optimum partial oxidation state resulted in ferromag-

netism while overoxidation or underoxidation both lead to the expected diamagnetic

124 behaviour.

51 The computational work of Kasai et al. shows that Ag surfaces exposed to O2 can result in metastable ferromagnetic interactions. Specifically, Ag(111) surfaces are

able to dissociate O2 only under certain conditions (e.g., vertical trajectory of O2 molecule as it approaches the Ag surface; and, the molecular interaction with certain

Ag atomic sites – for example, the top of an atom or fcc-hollow site). Once disso-

ciated onto Ag, O atoms (at a 0.5 monolayer coverage in their studies, equivalent

to a 2:1 Ag:O ratio) create magnetic states at the surface due to a combination of

superexchange interactions between neighboring O atoms via the Ag atom; and, direct

exchange interactions between the 4d orbitals of Ag and its nearest O 2p states.51

Moreover, these particular direct exchange interactions result in ferromagnetic be-

haviour, while interactions between the 2px orbital of the O atom and the 5s or- bital of the nearest Ag atom result in antiferromagnetic direct exchange interactions, although their contributions are small compared to the former. 51 The same group

studied oxygen diffusion (at a 0.5 monolayer coverage) on Ag(111). 52 They found

that in-plane diffusion (i.e., parallel to Ag surface) reduces the distance between O

atoms which could induce associate desorption as O2. Similarly, perpendicular dif- fusion (i.e., sub-surface O atoms) which results in O atoms gaining more electrons

from the Ag surface resulting in quenching of the ferromagnetic behaviour.52 It is

important to note that adsorbed O2 without dissociating (i.e.,asamolecule)atthe Ag surface does not induce a magnetic response but the event does occur. 51

In the current manuscript, we extend our studies to Ag in both NP and thin-

film forms. Ag is diamagnetic in the bulk with a molar susceptibility, χm, of –19.5 × 10−6 cm3 mol−1,77 where the negative sign indicates diamagnetic behaviour. For comparison, the most common Ag oxides Ag2O and AgO are both diamagnetic in

−6 3 −1 77 the bulk with χm of –134 and –19.6 × 10 cm mol , respectively. NPs and thin-

125 film synthesis was conducted by sputtering in the gas-phase thus avoiding capping

ligands and allowing for pristine surfaces to be analyzed. Our main findings are that

1) metallic silver NPs as-prepared are non-magnetic; 2) magnetism is indeed a surface

phenomenon, governed by surface oxidation; and, 3) we find no strong size dependence

on the magnetic properties. This is a significant finding that may yield insight into

the wide scatter of reported magnetic properties, e.g., saturation magnetization.73,156

5.2. Experimental

The specific experimental methodology for Ag nanomaterials including synthesis and

collection are found on Section 2.11

5.3. Results

5.3.1. Ag NPs

Ag NPs of varying diameters were prepared by modifying chamber conditions (see Ta-

ble 2.3 for parameters) while generating NPs. NPs with modal diameters of 3.3 ± 0.9,

5.0 ± 0.9, and 7.8 ± 1.3 nm were obtained as determined from in situ mass spec- trometry as depicted in Figure 5.1a. The experimental in situ mass spectrometry spectrum peak for the smallest NPs was fitted with a single Gaussian peak. However, the spectra for the remaining two Ag NP diameters of 5.0 ± 0.9 and 7.8 ± 1.3nmwere

fitted with two Gaussian peaks each as shown in Figure 5.2. This could indicate that the populations of these two NP diameters are comprised of two separate diameters each.

The source conditions presented in Table 2.3 were used to generate all the NPs studied in this work. Ag NPs were deposited for 10 min onto Si wafers in the collec- tion chamber. Following preparation, samples were stored in sealed containers under

Ar gas until SQUID measurements were performed. Figure 5.1b shows SQUID mag-

126 Figure 5.1: a. Ag NP sizes determined by in situ mass spectrometry. The responses have been shifted vertically, and the response for the smaller NPs has been multiplied by a factor of 2 for ease of visualization. Mean diameters are ()3.3± 0.9 nm, (×) 5.0 ± 0.9 nm, and ()7.8± 1.3 nm. Lines are Gaussian fits to the data. For the two bigger NPs, two Gaussian fits per peak were utilized (see Figure 5.2). b. SQUID measurements of Ag NPs at room temperature (300 K). Inset is closeup of the origin. Lines are guides to the eyes.

127 Figure 5.2: Ag NPs with a size distribution of (a)5.0± 0.9 and (b)7.8± 1.3 nm. Each panel consists of the experimental data (scattered symbol), two Gaussian peaks and overall fitting (solid lines).

128 netometry measurements on these NPs supported on their original Si wafer substrate.

A clear magnetic response is observed in the M versus μ0H loops for all NPs. Specif- ically, a ferromagnetic signature is observed as seen in the inset on Figure 5.1b which

corresponds to a closeup of the origin showing weak coercivity μ0Hc and remanent magnetization, which are indicative of the Ag NPs behaving as soft ferromagnets at

room temperature (300 K). Based solely on the M versus μ0H loops, there is no clear size dependency to the magnetic signal as all loops show a similar – albeit noisy –

−3 −1 magnetic saturation M s of ∼0.93 × 10 emu g , coercivity μ0Hc of ∼24 mT and

reduced remanent magnetization M r/M s of ∼0.16.

5.3.2. Ag films

In Figure 5.3, magnetization loops are shown for an ON-target film exfoliated from

its substrate using Kapton tape. Similarly to the Ag NPs, the film collected ON-

target shows a clear magnetic response. Moreover, closer inspection of the M vs H

loop near the origin (Figure 5.3 inset) shows a weak coercivity (μ0Hc =1.28mT)

and reduced remanent magnetization (Mr/M1T = 0.012), an indication that the Ag film behaves as a soft ferromagnet at room temperature. Also shown in Figure 5.3

are analogous data for a film collected OFF-target. In sharp contrast to the NPs

and ON-target film, the OFF-target sample shows no obvious hysteresis and the

film is thus considered diamagnetic.

It is important to consider that by exfoliating the film off its substrate, any possible role played by the substrate on the magnetic response of the film is negated. Figure 5.4 shows SQUID magnetometry measurements of a Ag film collected ON-target onto

NaCl and exfoliated with Kapton tape. The sample shows hysteresis which suggests that the observed ferromagnetism is independent of substrate as it rules out epitaxial driven formation or other such substrate-dependent processes.

To gain some insight into the dramatic different magnetic behaviours observed for

129 Figure 5.3: Room-temperature (300 K) M vs H loops for Ag thin films collected ON- and OFF-target. Inset shows the zoomed section near the origin.

OFF-target vs. ON-target films, AFM and SEM microscopies were used to study

the morphologies of the samples. In Figure 5.5, the topographies of the ON-target

and OFF-target samples are shown. The two samples were prepared simultane-

ously (eliminating extrinsic impurities) during a 5-min deposition onto Si. Despite

the identical source conditions, the SEM and AFM images show a drastic morphology

difference. As seen, the film collected OFF-target is relatively flat containing ∼65

nm diameter islands and a relatively homogeneous size distribution, typical of vapour

deposition of atomic Ag species onto the Si substrate which produces mirror surfaces

(Figures 5.5a and 5.5b). On the contrary, the film collected ON-target shows fea-

tures of ∼1 μm in size in both SEM and AFM (Figures 5.5c and 5.5d). Moreover, both Figure 5.5c-d show the presence of small structures of a few nanometers in size dispersed on the surface of the 1 μm structures not present in the other film.

Figure 5.6 shows the presence of NPs is characteristic of ON-target deposition, irrespective of deposition time. The difference in topographies does suggest that

130 Figure 5.4: Room-temperature (300 K) M vs H loops for a ON-target Ag thin film transferred to Kapton tape after direct deposition onto a NaCl substrate. Inset shows full applied field range. Main graph is the zoomed section near the origin. Lines are guides to the eye. the presence of nanodimensioned structures is pre-requisite for ferromagnetism to be observed. This hypothesis is consistent with the fact that the NP films also exhibited ferromagnetic behaviour, although varying the size of the nanoparticles in the 3 to 8 nm diameter range had no apparent effect.

5.3.3. Ag NPs: In Situ Oxidation Effects

In previous work, ferromagnetism of Cu NPs and nano-structured films was associ- ated with non-stoichiometric oxidation. 94 To explore this possibility, in situ MQCM magnetic measurements of nascent Ag NPs were made as a function of their expo- sure to oxygen. Representative data are presented in Figure 5.7. We have previously shown the responses of different magnetic materials during MQCM experiments.94

The general trends in MQCM experiments are as follows: ferromagnetic materials re- sult in a frequency increase (i.e., positive Δf MQCM) when exposed to a magnetic field,

131 Figure 5.5: SEM (top) and AFM topography (bottom) images of Ag thin films col- lected OFF-(a and b)andON-target (c and d). All scale bars measure 200 nm. while diamagnetic and weakly-paramagnetic materials result in a negative frequency change. As seen in Figure 5.7, when the mass of oxygen adsorbed to the NPs is zero

(i.e. nascent NPs), a small negative frequency change when exposed to a magnetic

field of 0.1 T is observed; this is indicative of a non-magnetic species.94 As oxygen

is adsorbed, there is a rapid increase in magnetic response that abruptly stops and

proceeds to slightly diminish with further oxidation. This magnetic response is per-

sistent even with heat treatments of the samples in air. Clearly, there is a correlation

between the magnetic properties of the Ag NPs and oxidation.

To explore the role of oxidation, 3.3 ± 0.9nmAgNPsandON-target thin films were deposited onto NaCl surfaces. Samples were prepared by depositing nanomate- rial for 30 min followed by their immediate storage in either Ar or O2 atmospheres for a period of 5-6 days. It is important to note that all samples were exposed to air

for a short period of time during sample retrieval from storage, and during sample

preparation and mounting for magnetometry measurements. Thus minor oxidation

is expected even for samples kept in Ar. To perform SQUID magnetometry measure-

132 Figure 5.6: SEM images of ON- and OFF-target films. (a) ON-target deposition for 1 min and (b) 15 min. (c) OFF-target deposition for ∼2hr.

133 Figure 5.7: MQCM experiment using () nascent Ag NPs with subsequent ()ox- idation with ambient air and () heat treatment. All measurements performed at room temperature. Line is a guide to the eye.

ment, NP and ON-target films were exfoliated from the NaCl surfaces with Kapton

tape. SQUID magnetometry data for NP samples are shown in Figure 5.8a. As seen,

the sample stored in O2 had a higher saturation magnetization M s than the sample stored in Ar gas.

In addition to higher saturation magnetization M s for Ag NP samples stored in

O2 compared to Ar gas, the remanent magnetization M r/M s and coercivity μ0Hc (lower inset of Figure 5.8a), show higher values for the Ar-kept sample compared

to O2-kept NPs. This data is found in Table 5.1. Data for the analogous SQUID measurements for ON-target Ag films are shown in Figure 5.8b and Table 5.1. The effect of oxidation on the ON-target Ag film samples was analogous to the NP samples with those stored in oxygen having higher saturation magnetization M s that those stored in Ar. In addition, Figure 5.8 suggests the presence of a weak magnetic

field during Ag deposition (as in the case for ON-target films but not NPs) does not

play a significant role in the occurrence of magnetism.

134 Figure 5.8: Room-temperature (300 K) SQUID magnetometry measurements of (a.) Ag NPs and (b.) thinfilmsstoredineitherAr()orO2 () atmospheres. Bottom insets in both (a.)and(b.) show a close up of the origin. Lines are a guide to the eye. Top insets show SEM images with a (a.) 100 nm and (a.) 500 nm bars.

135 M s M r/M s μ0Hc (10−3 emu g−1)(mT) NPs

Kept in O2 4.28 0.20 12.62 Kept in Ar 1.40 0.25 28.56 ON-target thin films

Kept in O2 1.45 0.03 0.95 Kept in Ar 0.81 0.08 8.64 Heat-treated ON-target thin films

As prepared kept in O2 1.29 0.02 1.15 75 ◦C for 2 hrs 1.06 0.05 7.73 150 ◦C for 2 hrs 0.94 0.10 12.39

Table 5.1: Magnetic properties of Ag nanomaterials.

To further explore the effect of oxidation, three Ag films collected ON-target and deposited onto Si were heat-treated in air (i.e., 75 and 150 ◦Cfor2hrseach) with one of them tested as-prepared. The series of SQUID measurements of heat treated films is shown in Figure 5.9. As seen, heating in air results in a decrease in

◦ saturation magnetization M s; however, subsequent heating (e.g., 150 C) causes less of a decrease than the prior heating (e.g.,75◦C). In this context, the SQUID data are consistent with the MQCM data presented above with oxidation resulting in a higher magnetic response than that observed in more fully oxidized samples. Another notable change in these samples is their coercivity μ0Hc and remanent magnetization

M r/M svalues (see Table 5.1).

5.3.4. Surface Analysis

To determine the impact of any surface contamination on the magnetic properties observed, EDXS analysis was performed on an ON-target Ag film resulting from a 15 minute deposition onto a NaCl substrate. Figure 5.10 shows EDXS elemental analysis, using two excitation beam energies (15 and 30 keV) that penetrate the film

136 Figure 5.9: Room-temperature (300 K) SQUID magnetometry measurements of heat-treated Ag films. Films were heat-treated in air at 75 ◦C() and 150 ◦C(). An as-prepared sample stored at room temperature under O2 ()isalsoshown.

to different extents (estimated at ∼500 nm and ∼1250 nm, respectively, using the

CASINO simulation software188). Except for a trace level peak observed at 1.24 keV

for the spectrum collected with 30 keV excitation (identified as Mg – a diamagnetic

element), only transitions assigned to Ag are observed. In addition, the ferromagnetic

behaviour of the film is clear from the M versus μ0H loop presented in Figure 5.4. This would indicate that the film does not contain any magnetic impurities (i.e.,Fe,

Co or Ni) that could induce the observed ferromagnetic behaviour. The noticeable

absence of oxygen shows the bulk of the material consists of metallic silver. This

result is consistent with the MQCM results of NPs behaving ferromagnetically only

after exposure to ambient air with subsequent oxidation (Figure 5.7). As EDXS has

limited surface sensitivity, XPS was used to analyze the surface of the deposited Ag

material.

To gain insight into oxidation states of Ag species that might contribute to the

ferromagnetism observed, XPS analysis was performed on all samples presented in

137 Figure 5.10: EDXS analysis of a Ag thin film transferred to Kapton tape after depo- sition onto NaCl substrate. Main graph shows an energy beam of 30 keV while inset shows a spectrum collected with a 15 keV energy beam. The zero peak present at 0 V in both spectra is an instrumental noise peak.

Figure 5.9. XPS surveys of these Ag thin films are shown in Figure 5.11.

Figure 5.11 supports the lack of magnetic impurities shown by EDXS. Further- more, XPS data showing the Ag 3d region of Ag films kept in O2 before and after sputtering, and heat-treated at 75 and 150 ◦C are presented in Figure 5.12a with their corresponding O 1s region in Figure 5.12b. For the 150 ◦C heat-treated Ag sample, the presence of AgO in the O 1s region (i.e., Figure 5.12b) is evident; however, peak

fitting in the Ag 3d region (i.e., Figure 5.12a) for this particular Ag species is unclear as seen by the minute dashed peaks which we interpret as AgO.

5.4. Discussion

5.4.1. Surface Species Characterization

In general, Ag does not oxidize at room temperature in O2; however, when exposed to moderate temperatures in the presence of O2, Ag dissociatively adsorbs O but does

138 Figure 5.11: XPS survey spectra and peak assignment for the sample kept in oxygen and heat-treated Ag films. Spectra show Ag as the major component with common atmospheric elements such as C and O.

139 Figure 5.12: XPS data showing the (a)Ag3d and the (b)O1s regions of Ag films ◦ kept in O2 before and after sputtering, and heat-treated at 75 and 150 C. Raw data are shown as scattered plots. All amplitudes are maximized for ease of visualization, thus are not comparable. a. Peak fitting and Shirley background are shown for the 3d 5/2 and 3d 3/2 peaks. The appearance of the plasmon peaks are indicative of Ag metallic interior. b. Peak fitting (solid and dotted lines) and Shirley background are shown for the O 1s region. C−OandC−O peak areas in the O 1s region match those corresponding to the C 1s region (latter data not shown). 140 189 not form bulk oxide. Moreover, O2 can also be adsorbed without dissociation. Ex- perimentally, it is known that ligand-free solution-synthesized Ag NPs stored as dry

powders form metal oxide at the surface when exposed to air for prolonged periods of

time.190 Thus, oxygen interacting with metal silver surfaces forms silver oxide over-

191 layers. The two most common Ag oxides are Ag2O and AgO. Thermodynamically,

Ag2O is the most stable oxide in air at room temperature among the possible Ag oxides. 189,192 In contrast, AgO is stable under ultra high vacuum conditions but once

116 exposed to air, its surface decomposes to Ag2O acting as an insulating oxide. The

+ 2− 116 + bonding in Ag2O is classified as ionic containing Ag and O ; electronically, Ag is diamagnetic. In contrast, AgO contains two non-equivalent Ag sites: a Ag+ coor-

dinated twofold by O atoms, and an almost square planar Ag3+.116,189 Experimental

neutron diffraction analysis of AgO confirmed its diamagnetic characteristic.189,193

Thus, both of these oxides are diamagnetic in the bulk. 189,193 Accordingly ferromag-

netism observed in our Ag films must be associated with non-stoichiometric oxides,

consistent with the observation that magnetic susceptibility is maximum when the

Ag films are partially oxidized as indicated by the MQCM data in Figure 5.7 and the

SQUID data in Figures 5.8 and 5.9.

QCM data allow us to calculate the average stoichiometry of Ag oxides present

on the surface of the films at the peak of magnetic susceptibility in the MQCM data

(Figure 5.7). This peak corresponds to an uptake of 0.7993 μg of oxygen by ∼97 μg

Ag NPs deposited onto the QCM crystal. This leads to a Ag:O atomic ratio of ∼8:1.

In order to calculate this ration, a few assumptions were made. For instance, it is

assumed that all NPs are 3.3 nm in diameter (based on average experimental results

of smaller NPs) and are perfect spheres. Then, it can be estimated that each NP

NP × 3 contains Nttl =1.50 10 Ag atoms (based on an atomic radius rAg of 144 pm).

The volume of the last atomic shell VOS on a NP of radius rNP can be estimated as

141 4πr3 4π V = NP − (r − 2r )3 (5.1) OS 3 3 NP Ag NP × 2 from which it can be estimated that there are Nsurface =6.58 10 surface Ag atoms / Ag NP. The total number of surface Ag atoms Nsurface in a 96.6 μg Ag sample (containing Ag Nttl )isthengivenas

NP Ag Nsurface · Ag × 17 Nsurface = NP Nttl =2.36 10 (5.2) Nttl Knowing this sample uptakes 0.7993 μg of oxygen from QCM measurements (i.e.

O × 16 Nttl =3.01 10 ) and assuming these O-atoms will react at the surface, the ratio of surface Ag atoms to O atoms is then

N Ag 2.36 × 1017 surface = =7.84 ∼ 8 (5.3) O × 16 Nttl 3.01 10 It must be noted that the ∼8:1 Ag:O atomic ratio is an overestimated value as the calculation assumes a film of NPs in monolayer arrangement in which all Ag atoms are surface atoms and available for oxidation. MQCM analysis during extended heat treatments with increasing oxygen uptake do not seem to induce an effect (either a positive or negative Δf MQCM) and the signal intensity remains stable. As seen in Figure 5.7, the greatest change in MQCM response by the Ag NPs is with initial exposure to oxygen and initial heat treatments. The lack of additional change in magnetic response with additional heating likely indicates that the additional mass being adsorbed is not manifesting itself in the formation of non-stoichiometric oxides but rather that the increase in temperature affords possibility of oxidation of intersti- tial Ag surface available in the film interior thus forming stoichiometric, non-magnetic oxides as expected withing the framework of a “normal” Ag lattice. In previous work the accessibility of oxygen to the interior of Cu NPs (a more readily oxidizing metal compared to Ag) was modeled and experimentally validated.173 That work is consis-

142 Figure 5.13: XPS literature comparison to our Ag 3d5/3 peak positioning. (a)Liter- ature average from the NIST database.196,197 (b) XPS values from Ref. 195. Labels marked with an * represent deconvoluted peak values. (c)Our experimental XPS val- ues. Each panel consists of metallic Ag (bottom of ordinate) and several oxide species. Broken lines represent the binding energy division (i.e., higher or lower) compared to metallic Ag.

tent with only the outer 1-2 layers of exposed Cu atoms being available for oxidation

at room temperature.

In general, higher binding energy shifts are observed in XPS analysis when metals

are compared to their respective oxides. This is believed to be the result of electroneg-

ativity differences between the metal and its cation, as the latter has a lower electron

density in the valence region which shields the nucleus resulting in an increased bind-

ing energy of the core electrons compared to the metal. 194 However, the opposite

trend has been historically observed for Ag and its bulk oxides194 as presented in Fig-

ure 5.13a. Possible factors responsible for such phenomenon are ionic charge shifts,

lattice potentials, and extra-atomic relaxation.170,194 Moreover, XPS peak assignment

forAganditsoxidesisnotstraightforwardasthe3d-region peak position shifts are small and often overlap making it difficult to determine the oxidation state. 189 For

example, the most common peak utilized for Ag oxidation assignment is the 3d 5/2 which has been experimentally found between 367.9 and 368.4 eV for metallic Ag;

195 from 367.6 to 368.5 eV for Ag2O; and, from 367.3 to 368.1 eV for AgO –allof which overlap with each other.

143 A comparison of the Ag 3d 3/2,5/2 peak positions to literature values was carried out and found in Figure 5.13. Each panel on Figure 5.13 consists of metallic silver

accompanied by several silver oxide species. In each case, metallic silver is considered

the point of reference (broken vertical lines in this figure). Figure 5.13a is the general

trend obtained from a literature average of values found on the NIST database.196,197

As seen, the binding energy values for all silver oxides are lower than that of metal-

lic silver.194 Ionic charge shifts, lattice potentials, and extra-atomic relaxation are possible explanations for this uncommon phenomenon.170,194

Figure 5.13b consists of a more recent study of commercially available Ag and

Ag oxide powders including the deconvolution of the 3d 5/2 Ag peak. For a better comparison to our data, Figure 5.13b values have been adjusted using 284.8 eV C

1s rather than their original 285 eV C 1s calibration value. This study195 shows the

Ag+ species always present at higher binding energies than metallic silver (except

3+ in a mechanical mixture of AgO and Ag2O powders) while all Ag instances are found at lower binding energies. Figure 5.13c shows this work’s experimental XPS

data. Our data tracks the recent study195 better than the accepted values196,197 with

Ag2O the most likely species at the surface. It is important to note that our samples are not pure oxides but rather a layered system with oxide species at the surface or

subsurface of a metal interior which affects the Ag 3d5/2 peak position. The XPS data shown in Figure 5.12a (also presented more clear in Figure 5.13c)

reveal a shift of the Ag 3d 5/2 peak to higher binding energy upon exposure to O2 (relative to the sputtered surface which is effectively pure Ag). At its simplest, this

shift is consistent with an oxidation process forming Ag+ where shielding electron

density (e.g., s electrons) have been decreased. From the literature, both Ag2O

+ and the Ag contribution of AgO are shown by Figure 5.13b to have 3d 5/2 binding energies shifted to higher energy compared to the binding energy of Ag. With high

144 temperatures (150 ◦C) heating, however, the peak in Figure 5.12a (top frame) is seen to shift back to lower binding energy (also presented in Figure 5.13c). This shift is consistent with the formation of AgO species as can be seen in Figure 5.13a and b.

In this context, higher temperatures drive conversion of the Ag nanostructures to non-magnetic AgO species, while lower temperatures favour formation of Ag(I)-oxide species, some of which contribute to the ferromagnetism observed. This is supported by Figure 5.12b where AgO is clearly seen at 150 ◦C but under no other condition.

The O 1s spectra in Figure 5.12b also show clearly that Ag2O is the dominant species

◦ at O2, 75 and 150 C conditions. However, peak fitting clearly shows that Ag2O and AgO cannot by themselves be responsible for the spectra observed and at least two other chemically distinct species must be present in order to fit the shape of the spectra observed. While sum of these might be attributable to adventitious surface contamination, it is equally possible that they may also be attributable to non-stoichiometric oxides.

A useful characteristic of Ag XPS spectra is the presence of loss peaks due to plasmons, which result from the collective excitation of the valence band electrons by the photoelectron and are indicative of a clean surface. For Ag, plasmon loss peaks are expected at binding energies of ∼372 and ∼378 for the Ag 3d 5/2 and 3d 3/2 peaks, respectively, ∼4 eV higher than their corresponding photoelectron peak.198 All of our

XPS spectra (Figure 5.12) contain plasmon peaks at ca. 370.6 and 376.6 eV indicative of Ag(0) most likely present in the interior of the sample (taking into account that

XPS analysis probes at a maximum depth of ∼10 nm) as we expect the oxidation to be a surface effect.94 Furthermore, the plasmon peak position is 3.6 eV on average from their respective 3d 5/2 peak (see Table 5.2), as expected. The presence of such peaks in all of our XPS spectra are indicative of a thin-layer surface oxidation since plasmon peaks are indicative of a clean surface.

145 ◦ ◦ Sputtered Kept in O2 75 C 150 C

3d 5/2 position (eV) 368.2 368.4 368.4 367.8 3d spin-orbit split (eV) 6.00 6.05 6.00 6.00 3d 5/2 FWHM (eV) 0.665 0.694 0.696 0.670 percent Ag(0) present (%) 100.0 81.5 88.5 70.7 percent Ag oxide present (%) 0.0 18.5 11.5 29.3 Δplasmon–3d 5/2 position (eV) 3.70 3.55 3.60 3.60

Table 5.2: 3d 5/2 XPS fitting analysis for Ag thin films from Figure 5.12

In addition to the 3d Ag XPS analysis, Figure 5.12b shows the XPS analysis

for the 1s O region of the SQUID-analyzed samples presented in Figure 5.9. The

non-heated sample kept in O2 shows3peaksinthe1s O region indicative of metal oxide as the major component and of minor adventitious C species such as ether and

ketone (i.e., C-O and C=O, respectively). The presence of a metal oxide peak in

the 1s O region is evidence of the presence of Ag2Ointhe3d Ag region allowing for the deconvolution into two peaks as presented in Figure 5.12a. The sputtering

procedure, which removed ∼50 nm of surface material, performed on the non-heated sample resulted in a metallic Ag interior, with a 0.23% ether total component in the 1s O region. This is consistent with the observation of the plasmon loss peaks.

◦ Comparing the O2-kept Ag film to the heat-treated at 75 C shows that the same chemical species, albeit at different concentrations, are present. For instance, the

ether:ketone ratio present in the Ag sample kept in O2 is 2.32:1 compared to 1.67:1 in the heat-treated sample. A third Ag sample heat-treated at 150 ◦Cshowsthe

presence of two metallic oxides in the 1s O region indicative of AgO and Ag2O. In general, the interaction of molecular oxygen with materials is complex and

classification of oxygen species is broad. At the surface of materials, including sil-

ver, oxygen states can be classified into four groups: weakly-adsorbed physisorbed

oxygen, molecular chemisorbed oxygen, dissociatively-adsorbed atomic oxygen, and

subsurface oxygen incorporated below the surface or into the bulk.199,200 Moreover,

146 subsurface oxygen does not prevent continuing O chemisorption from the gas phase at the Ag surface.199 Additionally, as a result of Ag multivalence nature (i.e.,Ag+ and

3+ Ag ) the possible formation of several sub-oxide species such as Ag2O3,AgO,Ag2O

201 and Ag3O4 is not surprising. However, evidence for the presence of such species was not observed.

5.4.2. Origin of ferromagnetism in Ag NPs

In this work, three types of Ag samples were produced: NPs, and ON- and OFF- target films. As revealed by SEM and AFM, the former two of these included nan- odimensioned structures (i.e., inset in Figure 5.8a for NPs; and, Figures 5.5c, 5.5d,

5.8b inset, 5.6a and 5.6d for ON-target films) and these were found to be ferromag- netic when partially oxidized. The OFF-target films, although chemically identical

(prepared under identical conditions), formed a near-mirror like surface at 5-minute deposition with a very smooth morphology (i.e., Figure 5.5a–b); these OFF-target samples were diamagnetic (Figure 5.3) and therefore nanostructure is important for ferromagnetism to be present. However, it is also clear that oxidation is prerequisite for the magnetism observed in these Ag samples.

SQUID magnetometry measurements for both air-exposed Ag NPs (Figure 5.8a) and ON-target thin films (Figure 5.8b) displayed ferromagnetic behaviour. How- ever, the observed magnetic signal is not entirely due to similar nanodimensionality

(insets in Figure 5.8). To probe size effects over the range afforded by our experi- mental method, three NP sizes were investigated. All three NPs sizes ranging from

∼2–10.5 nm (Figure 5.1a) displayed hysteresis loops characteristic of soft ferromag- netic behaviour (Figure 5.1b). More importantly, there is no discernible trend with respect to correlation between saturation magnetization M s and NP diameter. Simi- larly, the comparatively large (by comparison) nano-dimensioned structures prevalent in the ON-target films (see Figures 5.5c and d), also manifested a soft ferromag-

147 netic response (Figure 5.3). In combination, these results suggest that it is more the

presence of a roughened surface topography or a high surface-to-volume ratio that is

important for inducing ferromagnetism in Ag films than the presence of nanoparticles

with specific diameters.

To further probe the effect of oxidation, ON-target thin-films were heated in air to further the oxidation of the sample. In Figure 5.9, a decrease in saturation magnetization M s of the heat-treated films is observed compared to the one kept in oxygen (Table 5.1). In particular, the sample heated at 150 ◦Cshowsthelowest

saturation magnetization of 0.94 × 10−3 emu g−1 and is the only sample containing

AgO (Figure 5.13c). As such, further oxidation is not conducive to forming a more

magnetic material. These results are consistent with those observed for Ag NPs in the

MQCM data (Figure 5.7). The combined data suggests that partial oxidation is key to

the observed ferromagnetism while additional oxidation leads to a reduced saturation

magnetization. This optimal degree of surface oxidation may provide insight into why

reproducibility of these magnetic behaviours is challenging.73

MQCM measurements of freshly-synthesized 3.3 ± 0.9nmAgNPs,forexample,

did not yield a ferromagnetic signal (Figure 5.7). This result indicates that nanodi-

mensionality alone is not sufficient to impart a magnetic behaviour to Ag, indicating

magnetism is an extrinsic property of the material. Magnetism can be activated upon

exposure to trace amounts of oxygen to form non-stoichiometric Ag-oxides AgnOwith

n <8. Further oxidation is not conducive to forming a more magnetic material since

higher temperatures yield non-magnetic AgO as evidenced by the XPS data. The

combined data suggest that partial oxidation of nanodimensioned structures is key to

the observed ferromagnetism while additional oxidation leads to a reduced saturation

magnetization.

Consistent with these findings, we propose the following process by which Ag

148 Figure 5.14: Proposed oxidation stages of Ag from a fcc unit cell to cubic Ag2O. Both metallic Ag and Ag2O are nonmagnetic. Intermediate steps include exposure to low O2 concentrations with a potential single type of magnetic centre, and high O2 exposure with two potential magnetic centre types. can become magnetic when exposed to oxygen as illustrated in Figure 5.14. Pristine

Ag nanostructures are diamagnetic. When the surface of these Ag nanomaterials is exposed to a low-O2 atmosphere, a non-stoichiometric oxide layer is formed at the sur- face. Such surface oxidation has been predicted to lead to ferromagnetic centres, 51,52 which from our experiments are characterized by low saturation magnetization and higher coercivity. Initially, further oxidation leads to a higher magnetization presum- ably due to the further formation of magnetic centres. At a certain point, however, further oxidation starts to reduce the magnetization of the material as clearly shown in the SQUID data in Figures 5.8 and 5.9. This may occur due to the increased antiferromagnetic interactions with distant surface sites,51 or the formation of the diamagnetic stoichiometric silver oxides. We also observe a decrease in the coerciv- ity as the magnetization decreases. Assuming a common (and yet-to-be-determined) magnetization reversal process is all samples, this can be due to the larger moment which makes reversal easier upon applying a magnetic field, or a decrease in the effective anisotropy constant. The latter can be due to the increasingly symmetric environment as the sample is further oxidized.

149 5.5. Conclusion

We have shown the synthesis of highly pure Ag nanoscale materials. Both Ag NPs and thin films behave ferromagnetically at 300 K. The ferromagnetic behaviour is extrinsic to the material and is linked to the nonstoichiometric oxidation near the surface. The ferromagnetism exhibited does not show size dependency. This study sheds insight into the origin of the unexpected ferromagnetic characteristics of nanoscaled coinage metal materials.

150 Chapter 6

Preliminary studies on gold

6.1. Introduction

Gold was the first coinage metal in which unconventional magnetism was observed.

The initial report of ferromagnetic behaviour was discovered in 2.5 nm Au NPs coated

with a self-assembled monolayer (SAM) of poly-N-vinyl-2-pyrrolidone resulting in a

37 saturation magnetization M s of ∼1.8 emu/g. This was a striking discovery since

−6 3 Au is diamagnetic in the bulk with a molar susceptibility, χm, of –28 × 10 cm mol−1,77 where the negative sign indicates diamagnetic behaviour. Since SAMs are well-known to impart new surface capabilities to materials,26 for instance they can

provide solubility in a particular solvent, the focus in the field of unconventional

magnetism has been the use of diverse capping molecules to induce ferromagnetic

behaviour.

Several instances of Au nanomaterials modified at the surface with SAMs, in

the process inducing ferromagnetic behaviour, have been reported in the literature.

Table 6.1 shows a small literature sampling of this phenomenon in Au NPs with an

emphasis on the wide range of saturation magnetization, M s,values.

151 It is important to note that Table 6.1 encompasses a wide range of diameters, testing conditions, and capping ligands; however, the variability of several magnetic characteristics for Au NPs has been noted by others.44,73 For example, Gr´eget et al.203 have reported their own variability in saturation magnetization M s, coercivity μ0Hc, and remanent magnetization M r of Au NPs capped with dodecanethiol stating that reproduction of magnetic properties with the exact same values could not be achieved in their case despite following strict control over synthesis parameters. Figure 6.1 shows the variation of M s values with respect to Au NP diameters reported by Gr´eget et al.203

In contrast to SAM-modified Au NPs, ferromagnetic behaviour in ligand-free Au thin films has also been reported. 44 In this particular case, Au thin films were syn- thesized by cluster-deposition consisting of 2.6 ± 0.6 nm in size. Thin films of 28 and

3 175 nm thickness resulted in saturation magnetizations M s of 10 and 1 emu/cm , respectively. The lower M s value for the thicker film was attributed to the presence of additional Au material resulting in a transition towards bulk-like behaviour.44 Sup- porting these findings are theoretical studies of Au12 and Au13 clusters with calculated

204 magnetic moments of 4 and 5 μB per cluster, respectively. In this chapter, following our success with ferromagnetic behaviour of Cu and Ag nanomaterials as a result of non-stoichiometric oxidation, we extend our studies to

NP diameter Ms Conditions Capping Ligand Reference (nm) (emu/g) 1.4 0.4 300 K and 1 T DDT 38 1.9 ± 0.2 0.02 2.6 K and 7 T PAAHC 202 1.9 ± 0.2 4.8 300 K and 6 T DDT 40 2.5 1.8 4.2 K and 5 T PVP 37 2.5 0.1 300 K and 4 T THPC 36 5-30 0.04 300 K and 1 T none 36

Table 6.1: Small literature sampling of Au NPs behaving ferromagnetically. Cap- ping ligands: polyvinylpyrrolidone (PVP), polyallyl amine hydrochloride (PAAHC), dodecanethiol (DDT), tetrakis(hydroxymethyl)phosphonium chloride (THPC).

152 Figure 6.1: Literature saturation magnetization M s values from SQUID magnetom- etry measurements at 295 K of Au NPs modified at the surface with dodecanethiol. Adapted from Ref. 203

Au. It is important to note that bulk Au is inert to surface oxidation as a result of its endothermic oxygen chemisorption energy. 44 Additionally, theoretical studies of

205 Au2O3 and Au2O indicated that these oxides are diamagnetic.

6.2. Experimental

6.2.1. Synthesis of Au nanomaterials

A negative aspect of the gas-synthesis method utilized to prepare our nanomaterials, which did not play a role with either Cu or Ag, is the issue of cost of the metal target.

The approximate cost of Cu targets was $10.00, followed by Ag at $50.00, and Au at $1,500.00. The high cost of the Au source target material limited our ability to prepare samples and to broaden our studies. Additionally, the synthesis process in itself is a resource consuming procedure as not all sputtered material becomes a NP.

This is a common problem with top-down approaches during the synthesis of metal

153 nanomaterials.206 For our experiments, a single Au NP diameter was synthesized and

deposited on NaCl, while Au thin films were deposited onto Si(100) substrates.

6.2.2. Other procedures

All other experimental protocols have been previously described in Chapters 2, 4,

and 5.

6.3. Results and discussion

6.3.1. Au NPs and thin films

The conditions utilized to synthesize our Au NPs were based on those determined to

prepare the smallest Ag NPs. These conditions can be found on Table 6.2. Sizing of

the resultant Au NPs was performed by mass spectrometry as shown in Figure 6.2.

The resulting Au NP diameter was determined to be 3.9 ± 1.7 nm.

NP diameter Ar flow rate He flow rate IVDp (nm) (sccm) (sccm) (mA) (V) (cm) (Pa) 3.9 ± 1.7 50.5 80.7 254 313.4 6 54.7

Table 6.2: Conditions for the generation of Au NPs. I is the sputtering current, V is the sputtering voltage, D is the distance of the metal target to the aggregation chamber exit, p is the pressure in the aggregation chamber during deposition.

Both Au NPs and ON-target thin films samples were collected under the same conditions. Au NPs were utilized for MQCM procedures and for SQUID analysis.

Similarly, ON-target thin films were collected for SQUID analysis. Both types of nanomaterials were deposited for 30 min in 5-min intervals with 5-min rest in order to allow the target to cool down. A difference between NPs and thin film samples acquired for SQUID analysis was the collection substrate. Au NPs for SQUID analysis were collected onto NaCl substrates, while ON-target thin films were collected onto

Si(100) die substrates. After deposition, samples were stored in either Ar or O2

154 Figure 6.2: Mass spectrum of Au NPs size collected by in situ mass spectrometry. Their mean diameter was determined as 3.9 ± 1.7 nm. Line is Gaussian fit to the data. atmospheres until further analysis.

6.3.2. MQCM

MQCM experiments were performed as previously described in Sections 4.3.3 and

5.3.3. The data are shown in Figure 6.3. There are two types of MQCM data presented in Figure 6.3: the inset shows in situ measurements, while the remaining plot represents ex situ heat treatments of the same sample.

In situ MQCM of Au NP measurements (inset in Figure 6.3) starts with a blank

QCM crystal (a, ). As previously observed with Cu and Ag MQCM experiments, the observed negative Δf MQCM for the blank QCM crystal indicates a diamagnetic material. The following data points (b-e, ) represent several depositions of Au NPs.

Depositions were performed for 5-min before MQCM was performed. Each deposition

resulted in a mass gain of ∼7 μg per deposition. As shown in the inset, depositions b through d experience a positive Δf MQCM when compared to the blank MQCM

155 Figure 6.3: MQCM experiment using (, b-d) nascent Au NPs with subsequent (, f and g) oxidation with ambient air and () heat treatment. The blank QCM crystal without NPs is also plotted (, a). Inset represents in situ measurements. Other data represents ex situ measurements. All measurements performed at room temperature. See text for details.

– albeit still negative overall. Deposition e shows a decreased MQMC signal from the previous deposition. This last deposition resulted in an overall mass addition of ∼28 μg. The reduced Δf MQCM signal during the last deposition is similar to the

reduction in saturation magnetization M s in Au thin films with increasing thickness as mentioned in Section 6.1, and reported by others.203 Data points f and g ()

represent the first and second air exposure, respectively. Each air exposure consisted

in breaking the vacuum seal for 10 seconds before re-establishing vacuum conditions,

and performing MQCM technique. As depicted in the inset, there is a slight gain in

mass of oxygen by the Au NPs with each exposure as seen by the reported change in

position in the x-axis. The initial oxygen exposure (f) resulted in a larger negative

Δf MQCM when compared to its previous value of e. The second air exposure g resulted in a Δf MQCM increase to values similar to e. Ex situ MQCM experiments ( in Figure 6.3) consisted of heat-treatments rang-

156 M s M r/M s μ0Hc (10−3 emu g−1)(mT) Au NPs Kept in Ar 5.23 0.12 5.55 Kept in O2 2.60 0.11 10.37

Table 6.3: Magnetic properties of Au NP. ingfrom40to120◦C. As can be seen in Figure 6.3 exposure to heat resulted in large negative Δf MQCM values. It is clear from the overall MQCM data that extensive oxidation of non-bulk Au NPs by heat treatment does not create magnetic centers.

In fact, the MQCM experiment seems to indicate that none or partial oxidation re- sults in Δf MQCM changes resembling those of ferromagnetic materials (i.e.,apositive

Δf MQCM); however, these changes are small and require further analysis to validate them.

6.3.3. Superconducting quantum interference device magnetometry

6.3.3.1. Au NPs on NaCl

Au NPs deposited onto NaCl substrates were exfoliated using Kapton tape for SQUID magnetometry studies. Magnetometry measurements of NP samples stored in either

Ar or O2 atmospheres are presented in Figure 6.4. As seen, the sample stored in

Ar had a higher saturation magnetization M s than the sample stored in O2 gas. A comparison of the magnetic properties of these two Au samples from SQUID magne- tometry can be found in Table 6.3.

6.3.3.2. ON-target Au thin films on Si

Attempts to remove ON-target Au thin films deposited on Si(100) by exfoliation with Kapton tape for SQUID analysis failed. This was unusual as Au and Si lattice constants do not match and should result in an effortless exfoliation process. Our

first attempt consisted in SQUID testing the Au films together with the Si substrate;

157 Figure 6.4: Room-temperature (300 K) SQUID magnetometry measurements of Au NPs stored in either Ar or O2 atmospheres. Bottom inset shows a close up of the origin. Lines are a guide to the eye. however, the resulting MvsHloops contained a large number of unusable data points due to their high uncertainty values.

Our inability to exfoliate the films prompted us to examine these samples via powder X-ray diffraction (PXRD). Since our thin film collection was performed ON- target and deposition times were perhaps too long at 30 min, it was possible that the heat produced by the target during sputtering induced minor alloying of Au and Si at the interface. This was an unusual result as Au does not form silicides, instead forming an abrupt metal-Si interface.207

6.3.4. Powder X-ray diffraction

A PXRD spectrum of an ON-target collected Au thin film deposited onto Si(100) substrate is presented in Figure 6.5. Except for the single peak positioned at 69.01◦ corresponding to Si(100) substrate, all peaks in the diffraction pattern can be at- tributed to Au. The presence of several Au peaks and their corresponding planes is

158 indicative of polycrystalline Au thin films. The strongest peak at 38.06◦ indicates that the preferred orientation is Au(111). The presence of both film and substrate on the XRD diffraction pattern is inconclusive in terms of the presence or absence of alloy formation. A technique that could help understand the presence of alloying at the interface is grazing incidence XRD.208 Thus, our inability to exfoliate the Au film from the Si substrate is unknown.

Figure 6.5: X-ray diffraction pattern of Au thin films collected ON-target onto Si(100) substrate. The peak positioned at 69.01◦ belongs to the Si substrate. All other peaks belong to Au.

6.4. Conclusions

Our findings indicate that Au NPs are able to be behave ferromagnetically without the presence of a SAM modifying its surface. The higher saturation magnetization

M s from SQUID magnetometry of Au NPs stored in Ar compared to O2 atmospheres is indicative of the similar non-stoichiometric oxidation systems observed in ferromag- netic Cu and Ag; however, our data is currently insufficient for a definitive mechanism.

159 Additionally, the MQCM trend observed for Au is unlike the behaviour observed for

Cu and Ag in our earlier work. For those two metals, there is a significant decrease in Δf MQCM not observed for Au. A larger nanomaterial sampling and other charac- terization techniques are required to fully analyze the Au system.

160 Chapter 7

Final conclusions and future direction

7.1. Summary and final conclusions

The overall aim of this thesis was to gain a better understanding of the unconventional magnetism in coinage metal nanomaterials. The three factors under consideration in this work were: the elemental composition of the nanomaterials by synthesizing single-metal nanostructures of diamagnetic Cu, Ag, and Au; the type of nanostruc- ture prepared consisting of NPs or thin films; and, the surface modification of such nanomaterials. The latter two factors, namely nanodimensionality and surface mod- ification in the form of self-assembled monolayers, are believed to be the reason for the observed ferromagnetic behaviour.

In order to study the independent effects of nanodimensionality and surface mod- ifications, and their possible role in magnetic behaviour, the gas-phase synthesis was utilized to prepare metallic NPs and thin films. The main advantage of this sputtering procedure, compared to the most common solution synthesis, is the lack of capping ligands during the preparation process resulting in nanomaterials exhibiting pristine surfaces. Another synthetic advantage consisted in our ability to deposit nanostruc-

161 tures onto multiple substrates. This latter aspect facilitated the implementation of

optical studies as a preliminary component of this thesis. Finally, gas-phase synthe-

sis resulted in the preparation of highly pure sub-12 nm size NPs and thin films of

nanomaterials containing pristine surfaces. Material composition purity (i.e., absence

of ferromagnetic impurities) was determined by XPS and EDX, while NP diameters

were determined mainly by in-situ mass spectrometry complemented by AFM and

SEM imaging techniques.

The first experimental component of this work consisted in studying the optical properties of NPs and thin films by a home-assembled MOKE setup. These prelim- inary MOKE results indicated a possible magnetic behaviour of our nanomaterials.

Since MOKE is only a qualitative magnetic technique, these results warranted further studies with the more robust SQUID magnetometry.

SQUID magnetometry studies provided a quantitative and clear ferromagnetic behaviour in both NPs and thin films of Cu, Ag and, partially of Au. However, our sputtering apparatus required that our samples were exposed to ambient conditions during retrieval after synthesis. The well-known reactivity of nanomaterials and their handling in air provided grounds to question their pristine surface status. A lesser known magnetic technique known as MQCM, based on the mature quartz crystal microbalance technique, was implemented inside our sputtering system.

In situ MQCM measurements of NPs indicated that our pristine coinage metal

NPs in the sub-12 nm diameter range were diamagnetic. The onset of magnetism only appeared after the introduction of oxidation. Controlled MQCM and XPS studies further characterized this process as non-stoichiometric oxidation.

These findings showed that there is no size effect on the observed ferromagnetism as both NPs and thin films resulted in ferromagnetic behaviour. However, the surfaces of both nanomaterials were composed of micro- and nanodimensionality structures.

162 Thus we proposed that the fundamental mechanism for the induction of ferromagnetic

behaviour to the coinage metals is a combination of nanodimensionality dominated

by surface chemistry in the form of non-stoichiometric oxidation.

7.2. Future direction

The future direction for this work should concentrate on the interaction of oxygen

with the coinage metal atoms at the surface and sub-surface levels, as we have deter-

mined this to be the key source of ferromagnetism along with a nanodimensionality

component. There are a few experimental techniques that could assist narrowing the

source of magnetism and to provide improvements to our current characterization

and understanding.

7.2.1. X-ray magnetic circular dichroism

X-ray magnetic circular dichroism (XMCD) is a technique that, unlike SQUID mag-

netometry, not only determines if a material is magnetic or not, but is able to locate

the source of magnetization (i.e., it is atom specific), and to determine quantitatively how much of the total magnetic moment corresponds to the spin and to the orbital angular momenta of each elemental component (i.e., able to differentiate between the spin and orbital moments via sum rules).29,119 This technique utilizes circularly polar-

ized synchrotron radiation and will require transportation of our sample to a specific

location, for example to the Canadian Light Source in Saskatoon, Saskatchewan.

Ideally, a secondary vessel inside our collection chamber in the sputtering system

able to store our synthesized samples under vacuum conditions would be required

in order to maintain the samples’ surface pristine during transportation. An advan-

tage of XMCD is its ability to perform analysis under ambient air, not necessarily

under vacuum. Thus, samples could tested under several oxidation stages: in their

163 as-prepared state without oxidation, with ambient oxidation, and with modest heat-

treatment steps. This experiment would allow for monitoring of the magnetic be-

haviour during the different oxidation steps, and would present strong evidence on

the exact source of the magnetic moments.

7.2.2. In situ magneto-optical Kerr effect

The addition of viewing windows on the collection chamber in the sputtering system

to allow laser probing of the sample, along with an electromagnet could lead to in

situ magneto-optic Kerr effect (MOKE) measurements. This setup would allow for

a more accepted magnetic experimental technique when compared to MQCM since

both are qualitative methods. Moreover, a proper MOKE setup would result in data

consisting of MvsH loops which is one of the accepted forms of magnetic data plots

(i.e., when compared to MQCM plots).

Alternatively, following with the transportation vessel mentioned in the XMCD

section, a chamber fitted with viewing windows able to keep vacuum could be utilized

for MOKE measurements – although technically not in-situ but in vacuo MOKE if

this variation is adopted. Additionally, MOKE can be performed under ambient

conditions which lends itself to the same experimental testing mentioned for XMCD.

7.2.3. Reduction of nanomaterials

In this thesis, our results indicated that non-stoichiometric oxidation of nanomaterials

at the surface led to ferromagnetic behaviour. It would be interesting to expose these

materials to a reducing atmosphere, for example under N or H2 gas, and observe their magnetic signature. Ideally, if this experiment was performed over several cycles,

oxidation would lead to a ferromagnetic response, while reduction would lead to a

disappearance of any magnetic behaviour. It must be noted that the temperatures

at which these experiments would be performed should be mild temperatures in an

164 attempt to minimize topographical changes which could affect the data. Magnetic

measurements for this experiment could be perform with either XMCD or MOKE,

or both.

7.2.4. Solid-state 109Ag nuclear magnetic resonance

In general, nuclear magnetic resonance (NMR) is a spectroscopic technique utilized to determine purity and molecular structure of samples. Specific to our Ag studies, solid- state 109Ag NMR spectroscopy is a technique that, unlike solution NMR, presents

well-defined measurements of Ag complexes.209 Specifically, solid-state 109Ag NMR

has allowed to differentiate silver oxide states in oxide mixtures.210 This technique

could complement our already performed XPS studies. The disadvantage is that

109Ag NMR does not offer the flexibility that XPS offers in terms of analyzing all three coinage metals utilized in this thesis.

7.2.5. Magnetic signal longevity

An interesting aspect that could be performed without the need of new instrumenta- tion is the study of magnetic signature longevity. According to our data, the source of magnetism is non-stoichiometric oxidation at the surface, which lends itself to long term monitoring studies of neatly-kept samples specific for such studies. For example, by exfoliating a metallic thin film from its substrate using Kapton tape, and ensuring the sample is sealed and protected from ambient exposure (i.e., avoiding further oxi- dation and potentially affecting its signal intensity), the same sample could be tested with SQUID magnetometry monitoring its magnetic saturation multiple times during an extended period of time.

165 7.2.6. Surface rearrangement

During our experiments with Cu and Ag, we observed the presence of smaller nanos- tructures on the microstructures of thin-films. Since these nanostructures were imaged with AFM and SEM after exposure to oxygen, the origin of these nanostructures is not known. For instance, it would interesting to determine if these nanostructures are the result of morphology changes induced by oxidation, or if these shapes are the result of the sputtering process. To accomplish this proposed study, metallic samples would have to be imaged before and after oxygen exposure. Again, a transportation vessel able to keep the nanomaterial under vacuum as well as a high-vacuum AFM instrument would be required for imaging.

7.2.7. Sharper NP diameters

During NP synthesis, in situ mass spectrometry was utilized to monitor sizing rather than to use it as a filter (i.e., only allowing a specific size to reach the substrate).

Although our NP distributions per specific condition can be considered narrow distri- butions based on the values of our standard deviations, better control on size provides

NPs with closer characterization values that could result in less discrepancy in mag- netic characteristics (e.g., magnetic saturation, magnetic coercivity, and remanent magnetization) as those reported by others.44,73 The disadvantage of this proposed experimental work is the amount of time and material that would be required in order to prepare appropriate dense samples for analysis.

7.2.8. Other synthetic methods

Although sputtering was extremely important to our approach for the creation of pristine surfaces, perhaps another method could be use in order to create similar or identical clean surfaces without so much waste during the top-down approach. As

166 determined by our preliminary Au findings, the top-down is not conducive for high output research. Ideally, a synthetic method that favours a specific packing phase such could lead to comparison studies, for example, the magnetic characteristics of a metallic fcc thin film compared to a bcc phase. Provided that the atomic packing efficiency is different, access to these metallic atoms by the diatomic O2 molecule could lead to a better understanding of non-stoichiometric oxidation, and thus its magnetic behaviour. For instance, a fcc surface is atomically more dense (atomic packing factor of 0.74) resulting in oxygen access to subsurface atoms almost impossible compared to a bcc surface with an atomic packing factor of 0.68.211 This idea was based on the computational work of Kasai et al.51 which predicts certain Ag facets oxidize easier than others due to atomic accessibility and packing.

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Appendix A

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