Alternatives to Organic Acid Surface Modification of Zno for Excitonic Photovoltaics

ALTERNATIVES TO ORGANIC ACID SURFACE MODIFICATION OF ZNO FOR EXCITONIC PHOTOVOLTAICS

by Thomas M. Brenner A thesis submitted to the Faculty and the Board of Trustees of the Colorado School of Mines in partial fulﬁllment of the requirements for the degree of Doctor of Philosophy (Applied Physics).

Golden, Colorado Date

Signed: Thomas M. Brenner

Signed: Prof. Reuben T. Collins Thesis Advisor

Signed: Prof. Thomas E. Furtak Thesis Advisor

Golden, Colorado Date

Signed: Thomas E. Furtak Professor and Head Department of Physics

ii ABSTRACT

Surface modification of metal oxides with molecular monolayers is an effective strategy for tuning interface properties in excitonic devices employing metal oxides as charge accepting and transport layers. The most commonly used attachment chemistries are acid/base reactions employing organic acids. The use of acid/base chemistries has presented a problem for one of the most commonly used and promising metal oxides in excitonic devices, zinc oxide (ZnO). ZnO is easily etched by even weak organic acids, leading to non-ideal monolayers and the accumulation of surface complexes during etching, which is particularly problematic for ZnO-based dye sensitized solar cells (DSSCs). Two ways to address this issue have been explored. The first approach is to employ a triethoxysilane (TES)-based covalent attachment scheme instead of an acid/base reaction for attaching modifier molecules. We demonstrate that dipolar mixed monolayers of phenyltriethoxysilane-based molecules tune the work function of ZnO and the performance of bulk heterojunction photovoltaic devices containing modified ZnO layers. This indicates these modifiers are effective for tuning interfacial electronic structure.

The second approach is to investigate Zn1-xMgxO (ZnMgO) alloys in order to produce a more etch resistant material with similar electronic properties to ZnO. These alloys, when exposed to the prototypical modiﬁer benzoic acid (BA), demonstrate a steady-state, macroscopic etch rate that decreases up to an order of magnitude (at 20% Mg) compared to ZnO. Infrared spectroscopic characterization of BA-modiﬁed ZnMgO indicates a monolayer of BA attaches to the ZnMgO surface nearly instantaneously and remains throughout etching. These results suggest that ZnMgO is a promising alternative material that may alleviate some of the problems with ZnO etching. However, for applications of this material as a substrate for dye sensitization, the initial etch rate, and not the steady-state rate, is really the quantity of interest. We investigated the initial etch rate of ZnMgO exposed to N3 dye (cis-

iii bis(isothiocyanato)bis(2,2’-bipyridyl-4,4’-dicarboxylato)-ruthenium(II)). We ﬁnd the initial etch rate of ZnMgO increases with Mg content, in contrast to the steady-state etch rates observed for BA-treated ZnMgO. We also ﬁnd that the primary products of etching are Zn-carboxylate products. From these results we propose a mechanism for the observed etch resistance.

iv TABLE OF CONTENTS

ABSTRACT ...... iii

LIST OF FIGURES ...... viii

LIST OF TABLES ...... xii

LIST OF SYMBOLS ...... xiii

LIST OF ABBREVIATIONS ...... xvi

ACKNOWLEDGMENTS ...... xviii

CHAPTER 1 METAL OXIDE SEMICONDUCTORS IN EXCITONIC PHOTOVOLTAICS ...... 1

1.1 Metal Oxide Semiconductors ...... 1

1.2 Organic and Fullerene Semiconductors ...... 4

1.3 General Photovoltaic Device Physics ...... 10

1.4 Excitonic Solar Cell Device Physics ...... 17

1.5 Metal Oxide/Organic Interfaces ...... 21

1.6 Monolayer Modiﬁcation of Metal Oxide Surfaces and Interfaces ...... 25

1.7 Acid Dissolution of Metal Oxide Semiconductors ...... 27

1.8 Thesis Organization ...... 29

CHAPTER 2 SAMPLE PREPARATION AND CHARACTERIZATION TECHNIQUES ...... 31

2.1 Production of Zn1-xMgxO Thin Films ...... 31

2.2 Triethoxysilane Modiﬁcation of ZnO ...... 32

2.3 Carboxylic Acid Modiﬁcation of Zn1-xMgxO ...... 33

v 2.4 Fabrication of Bulk Heterojunction Solar Cells ...... 34

2.5 UV-Vis Absorption Spectroscopy ...... 35

2.6 Infrared Absorption Spectroscopy ...... 36

2.7 Kelvin Probe Surface Potential Measurements ...... 39

2.8 Tapping Mode Atomic Force Microscopy ...... 42

2.9 X-Ray Photoelectron Spectroscopy ...... 45

2.10 Contact Angle Goniometry ...... 48

2.11 Grazing Incidence X-Ray Diﬀraction ...... 49

2.12 Photoluminescence Spectroscopy ...... 51

CHAPTER 3 TUNING ZINC OXIDE/ORGANIC ENERGY LEVEL ALIGNMENT USING MIXED TRIETHOXYSILANE MONOLAYERS ...... 53

3.1 Introduction ...... 54

3.2 Experimental ...... 58

3.3 Results and Discussion ...... 61

3.4 Conclusions ...... 71

3.5 Acknowledgements ...... 71

CHAPTER 4 ETCH-RESISTANT ZN1-XMGXO ALLOYS: AN ALTERNATIVE TO ZNO FOR CARBOXYLIC ACID SURFACE MODIFICATION . . . . . 72

4.1 Introduction ...... 73

4.2 Experimental ...... 76

4.3 Results and Discussion ...... 78

4.4 Conclusions ...... 92

4.5 Acknowledgements ...... 92

vi CHAPTER 5 EXPLORING THE MECHANISM OF ZN1-XMGXO ETCH RESISTANCE THROUGH DYE SENSITIZATION ...... 94

5.1 Introduction ...... 96

5.2 Experimental Methodology ...... 99

5.3 Results and Discussion ...... 104

5.4 Conclusions ...... 116

5.5 Acknowledgements ...... 117

CHAPTER 6 CONCLUSIONS ...... 119

6.1 Project Conclusions ...... 119

6.2 Future Project Suggestions ...... 122

REFERENCES CITED ...... 125

APPENDIX A - FURTHER EXPLANATION OF METHODOLOGY ...... 140

A.1 Infrared Active and Inactive Modes: An Example ...... 140

A.2 Fourier Transform Infrared Spectrometer Design ...... 141

A.3 Simpliﬁed Schematic of the Kelvin Probe ...... 142

A.4 TM-AFM Feedback Loop Schematic ...... 144

APPENDIX B - SUPPORTING INFORMATION FOR CHAPTER 3 ...... 145

B.1 Calculation of Surface Proportion of 4CPTES and PTES from Infrared Spectrum ...... 145

B.2 Water Contact Angle Measurements ...... 148

B.3 Atomic Force Microscopy Measurements ...... 148

B.4 Dark J-V Curves ...... 150

B.5 Eﬀect of Light Soaking ...... 150

APPENDIX C - PERMISSIONS ...... 152

vii LIST OF FIGURES

Figure 1.1 Hexagonal wurtzite structure of ZnO ...... 3

Figure 1.2 Atomic force microscopy height images of Zn1-xMgxO ﬁlms produced by a sol gel process ...... 5

Figure 1.3 UV-Vis spectra of the thin ﬁlms of Zn1-xMgxO studied in this thesis . . . . 6

Figure 1.4 Examples of common electro-active organic polymers and small molecules . 7

Figure 1.5 Illustration of Fermi level equilibration in junctions of electronic materials ...... 11

Figure 1.6 Example of a current density - voltage curve for a solar cell ...... 14

Figure 1.7 Equivalent circuit model of a solar cell ...... 17

Figure 1.8 The operational steps of excitonic solar cells ...... 18

Figure 1.9 Examples of excitonic photovoltaic device architectures ...... 19

Figure 1.10 Examples illustrating issues with metal oxide/organic semiconductor interfaces ...... 22

Figure 1.11 Electronic structure of a metal oxide and an organic in isolation and in contact ...... 23

Figure 2.1 UV-Vis absorbance spectra of materials used in this thesis ...... 37

Figure 2.2 Diagram of the working principle of the PM-IRRAS technique ...... 40

Figure 2.3 Basic working principle of Kelvin probe surface potential measurements . 41

Figure 2.4 Behavior of the tip in contact mode atomic force microscopy and tapping mode AFM ...... 43

Figure 2.5 Phase shift between AFM tip amplitude and driving force during TM-AFM ...... 44

Figure 2.6 TM-AFM height and phase images showing the relationship between topography, sample inhomogeneity, and phase ...... 45

viii Figure 2.7 X-ray photoelectron spectroscopy experimental setup ...... 47

Figure 2.8 Contact angle measurement setup and examples ...... 48

Figure 2.9 Experimental setup of grazing incidence X-ray diﬀraction ...... 50

Figure 2.10 Illustration of photoluminescence spectroscopy experiment setup and spectrum of N3 dye ...... 52

Figure 3.1 Energy level alignment at an ideal metal oxide/organic interface can be tuned by introducing a dipolar monolayer at the interface ...... 56

Figure 3.2 PM-IRRAS measurements of ZnO ﬁlms treated with PTES and 4CPTES in diﬀerent proportions ...... 63

Figure 3.3 Relative work function of treated ZnO as a function of 4CPTES and PTES mole fraction ...... 65

Figure 3.4 Representative light J-V measurements of IBHJ photovoltaic devices containing mixed monolayer modiﬁed ZnO ...... 67

Figure 3.5 Plot of open-circuit voltage of devices against relative work function of treated ZnO ﬁlms ...... 70

Figure 4.1 Bandgaps of ZnMgO ﬁlms as a function of Mg content ...... 79

Figure 4.2 Background-subtracted grazing incidence X-ray diﬀraction spectra of Zn1-xMgxO...... 80

Figure 4.3 PM-IRRAS infrared absorption spectra of benzoic acid-treated ZnMgO ﬁlms showing the carboxyl stretch region ...... 83

− Figure 4.4 Peak ﬁts to the dominant νasym(CO2 ) feature of the infrared spectra of ZnMgO ﬁlms soaked in benzoic acid for 1 hr ...... 84

− Figure 4.5 Ratio of the integrated intensity of the νasym(CO2 ) modes observed on benzoic acid treated ZnMgO ...... 86

− -1 Figure 4.6 Variation of the νasym(CO2 ) modes at 1532 and 1569 cm in ZnO samples as a function of exposure time to benzoic acid ...... 88

-1 − Figure 4.7 Plot of integrated intensity of 1550 and 1571 cm νasym(CO2 ) modes and their sum as a function of bandgap ...... 89

ix Figure 4.8 Relationship between amount of unreacted acetate in ZnMgO ﬁlms and amount of attached benzoic acid ...... 90

Figure 5.1 Molar absorptivity of N3 dye in ethanol solution ...... 101

Figure 5.2 Sensitivity factors of the orbitals measured during XPS as a function of their binding energy ...... 103

Figure 5.3 Bandgaps for ZnMgO samples with 0 - 20% Mg used in Chap. 5 . . . . 104

Figure 5.4 UV-Vis and PL spectra of ZnMgO treated with N3 dye for 1 hr . . . . . 105

Figure 5.5 Normalized PL spectra of ZnMgO treated with N3 dye for 1 hr . . . . . 106

Figure 5.6 Surface coverage and PL integrated intensity of ZnMgO treated with N3 for 1 hr ...... 107

Figure 5.7 PM-IRRAS Spectra of ZnMgO exposed to N3 dye for 3 hr and 24 hr . . 109

Figure 5.8 XPS composition measurements of ZnO and Zn0.8Mg0.2O before and after treatment with 2 mM benzoic acid for 30 min ...... 111

Figure 5.9 XPS composition measurements of Zn and Mg before and after treatment of Zn0.8Mg0.2O by BA for 30 min ...... 112

Figure 5.10 UV-Vis absorbance spectra of ZnMgO ﬁlms sensitized with N3 after undergoing a mineral acid pre-etch ...... 113

Figure 5.11 Surface coverage as a function of total exposure time to N3 dye for mineral acid pre-etched ZnMgO samples ...... 114

Figure 5.12 N3 surface coverage of a ZnMgO sample exposed to N3 dye for 24 hr and a sample exposed to N3 for 2 hr after a pre-etch ...... 115

Figure A.1 Examples of infrared inactive and active modes of carbon dioxide. . . . 140

Figure A.2 Experimental setup of an FTIR spectrometer ...... 142

Figure A.3 Schematic of Kelvin probe circuit ...... 143

Figure A.4 Diagram demonstrating the feedback loop that controls the AFM microscope ...... 144

Figure B.1 Diagram demonstrating calculation of IR integrated intensities contributed by PTES and 4CPTES ...... 146

x Figure B.2 AFM phase images of monolayer treated ZnO surfaces...... 149

Figure B.3 J-V curves of monolayer treated ZnO P3HT:PCBM devices in the dark 150

Figure B.4 J-V curves of 4CPTES-treated ZnO devices before and after light soaking ...... 151

xi LIST OF TABLES

Table 1.1 Bandgaps of Zn1-xMgxO ﬁlms calculated from UV-Vis spectra ...... 5

Table 3.1 Characteristics of inverted bulk heterojunction devices with monolayer modiﬁed ZnO electron contacts ...... 66

Table 4.1 Etch rate of ZnMgO ﬁlms exposed to 2 mM benzoic acid in hexane . . . . 81

Table 4.2 Peaks identiﬁed in ﬁts of the IR spectra of BA treated ZnMgO ...... 84

Table B.1 Water contact angle measurements of monolayer treated surfaces . . . . . 148

Table B.2 RMS roughness determined from tapping mode AFM topographic scans of TES treated ZnO surfaces ...... 148

Table B.3 Shunt resistances for monolayer treated ZnO P3HT:PCBM devices in the dark ...... 151

xii LIST OF SYMBOLS

Relative dielectric constant, also molar absorptivity ...... ε

Boltzmann constant ...... k

Temperature ...... T

Fermi energy in semiconductor physics ...... Ef

Electric potential ...... Φ

Built-in potential ...... Vbi

Electron quasi-Fermi energy ...... Ef,n

Hole quasi-Fermi energy ...... Ef,p

Conduction state energy in inorganic and organic semiconductors ...... Ec

Valence state energy in inorganic and organic semiconductors ...... Ev

Electron volume concentration in conduction states ...... n

Hole volume concentration in valence states ...... p

Density of states of the lowest energy conduction states ...... Nc

Density of states of the highest energy valence states ...... Nv

Current density ...... J

Fundamental unit of charge ...... e

Hole mobility ...... µp

Electron mobility ...... µn

Gradient operator ...... ∇

Voltage (equivalent to electric potential) ...... V

xiii Short circuit current ...... Jsc

Open circuit voltage ...... V oc

Power generated by solar cell ...... P

Solar cell eﬃciency ...... η

Electrical current or light intensity ...... I

Fill factor ...... FF

Shunt resistance ...... RSH

Series resistance ...... RS

Fermi energy ...... Ef

Work function ...... φ

Energy diﬀerence between the LUMO level of an organic and the conduction band of a metal oxide ...... ELC

Hydrogen ion otherwise known as a proton ...... H+

Transmittance or transmittivity (equivalent) ...... T

Absorptivity, not to be confused with absorbance ...... Abs

Reﬂectivity ...... R

Wavelength ...... λ

Absorbance ...... A

Absorption coeﬃcient ...... α

Path length through sample ...... l

Molar concentration ...... c

Semiconductor bandgap ...... Eg

Indicates the change in a particular quantity ...... ∆

xiv Energy or electric ﬁeld (context-dependent) ...... E

Phase angle between tip amplitude and driving force in tapping mode atomic force microscopy ...... θ

Angular frequency of tapping mode atomic force microscopy tip ...... ω

Tip oscillation amplitude in tapping mode atomic force microscopy ...... A

Tip quality factor in tapping mode atomic force microscopy ...... Q

Tip spring constant in tapping mode atomic force microscopy ...... k

Inelastic mean free path of photoelectron in X-ray photoelectron spectroscopy . . . λIMFP

Contact angle ...... θc

X-ray angle of incidence in grazing incidence X-ray diﬀraction ...... δ

Scattering angle in X-ray diﬀraction ...... 2θ

Acid dissociation constant ...... pKa

Symmetric (sym) or asymmetric (asym) stretch of the carboxylate group - − (CO2 )...... νsym/asym(CO2 )

Water dissociation constant ...... k−w

xv LIST OF ABBREVIATIONS

Transparent Conducting Oxide ...... TCO

Indium Tin Oxide ...... ITO

Fluorine-doped Tin Oxide ...... FTO

Current density - voltage curve ...... J-V

Maximum power point of solar cell ...... MPP

Dye sensitized solar cell ...... DSSC

Highest occupied molecular orbital ...... HOMO

Lowest unoccupied molecular orbital ...... LUMO

Point of zero charge ...... pzc

Isoelectric point ...... IEP poly-3-hexylthiophene (semiconducting polymer) ...... P3HT phenyl-C61-butyric acid methyl ester (fullerene derivative) ...... PCBM poly(3,4-ethylenedioxythiophene):poly(styrenesulphonate) (conductive polymer blend) ...... PEDOT:PSS

Metal-to-ligand charge transfer ...... MLCT

Fourier transform infrared spectroscopy ...... FTIR

Polarization modulated infrared absorption spectroscopy ...... PM-IRRAS

Atomic force microscopy ...... AFM

Tapping mode atomic force microscopy ...... TM-AFM

Octadecanethiol ...... ODT

xvi X-ray photoelectron spectroscopy ...... XPS

Kinetic energy ...... KE

Binding energy ...... BE

Grazing incidence X-ray diﬀraction ...... GIXRD

Photoluminescence ...... PL

Bulk heterojunction ...... BHJ

Titanium dioxide ...... TiO2

Work function ...... WF

Monolayer ...... ML

Inverted bulk heterojunction ...... IBHJ

Shorthand for Zn1-xMgxO ...... ZnMgO

Benzoic acid ...... BA

Full width half maximum ...... FWHM

Incident photon-to-current conversion eﬃciency ...... IPCE

xvii ACKNOWLEDGMENTS

There are quite a few people who’ve helped me through the doctoral process who I need to thank. I’d like to thank my family for their support and encouragement throughout this process. I’m also forever grateful to Dayna Jacob for her companionship, encouragement, and patience during the years I’ve spent doing the work of this thesis. I’m also extremely grateful to my advisers, Reuben Collins and Thomas Furtak, for the advice, guidance, ideas, and mentorship they’ve given me. I’d also like to thank Dana Olson who has been a valuable mentor. A number of my fellow students have been instrumental in my learning and in the ac- complishment of this work. I’d like to thank Darick Baker for his guidance during my early years of graduate school. I’d like to thank Gang Chen for his camaraderie and collaboration throughout our graduate years together. Without Gang’s contributions, this thesis wouldn’t be what it is today. I’d also like to thank Thomas Flores, Erich Meinig, Paul Ndione, and Xerxes Steirer for their eﬀorts on the projects in this thesis. Again, without them, this thesis wouldn’t be what it is. I’d also like to thank all the friends I’ve made along this journey for their kind words, conversation, and companionship. Finally, I’d like to thank the members of my committee: Reuben, Thomas, Dana, and Mark Lusk, Brian Gorman, and Stephen Boyes. Their input and ideas during my thesis proposal, research, and defense have been extremely valuable.

xviii CHAPTER 1 METAL OXIDE SEMICONDUCTORS IN EXCITONIC PHOTOVOLTAICS

The term ’excitonic photovoltaics’ refers to solar cells in which light absorption generates excitons, bound electron-hole pairs, instead of free carriers. Excitonic photovoltaic devices have been around for about 25 years, beginning with the invention of the dye sensitized solar cell in the 1990s, which took advantage of an ultrafast charge transfer from a light absorbing dye to a transparent metal oxide semicondutor in order to generate photocurrents and photovoltages. While the efficiency of these devices quickly saturated, new devices appeared that involved all-soft matter active layers composed of semiconducting polymers or smaller organic chromophores and fullerenes, referred to as organic photovoltaics. These devices have quickly risen in efficiency since the discovery of the bulk heterojunction active layer morphology in the early 2000’s.[1] Companies and research institutions are reporting efficiencies above 12% and are designing multijunction organic devices that may surpass single junction devices in efficiency. Dye sensitized solar cells have been making a comeback recently as well, after more than a decade of stagnation. A number of different innovations have allowed the world record to grow to 12.3%.[2] The metal oxide materials and molecular attachment schemes discussed in this thesis have already played a large role in the progress of excitonic photovoltaic devices, and likely will continue to do so in the future. In this first chapter, an introduction to the materials and physics of excitonic solar cells, and the role that metal oxides play in them, will be provided.

1.1 Metal Oxide Semiconductors

Metal oxide semiconductors form a group of electronically useful materials that are often high-bandgap, which makes them transparent through most of the solar spectrum, and often ’intrinsically’ doped by defects in their structure, making them relatively good conductors suited to transporting carriers without further processing. The ability to dope intrinsically

1 or extrinsically to achieve suﬃcient conductivity for device applications without signiﬁcant free carrier absorption in the visible leads to the term transparent conducting oxide (TCO).

Notable metal oxide examples include zinc oxide (ZnO), titanium dioxide (TiO2), nickel oxide

(NiO), molybdenum trioxide (MoO3), tungsten trioxide (WO3), and tin oxide (SnO2). These materials haven proven to be extremely valuable in excitonic photovoltaics precisely for the reasons just mentioned. Many of them can also be deposited through cheap, solution-based processes, which are an added benefit. Extrinsic degenerate doping can be used to create transparent conducting oxides such as indium tin oxide (ITO), fluorine doped tin oxide (FTO), and Al or Ga doped ZnO that have sufficient conductivity to accomplish lateral transport of carriers without significant resistive losses. These materials are often used as transparent conducting contacts in optoelectronic devices. ZnO features prominently in this thesis. It has a hexagonal wurtzite crystal structure shown in Figure 1.1. ZnO has a bandgap of 3.2 eV. In the thin film form used in optoelectronic devices, ZnO has n-type conductivity that has been attributed to nonstoichiometric defects in its structure such as interstitial Zn atoms or oxygen vacancies that generate free electrons.[3] This origin of conductivity has been under some debate recently.[4] The processing conditions (such as annealing in oxygen or in the presence of hydrogen) can have a substantial influence on the conductivity of ZnO and are necessary to consider when comparing ZnO prepared in different ways.[3, 4] It is also relatively easy to form nanostructures of ZnO, compared to other metal oxides and this is one of its major advantages.[5, 6] ZnO is also a piezoelectric material.[7] It is particularly susceptible to dissolution by even weak organic acids, one of the major subjects of this thesis.[8, 9] The variety of properties and applications of ZnO that have been studied are too broad to be discussed here, but several reviews are available.[3–6, 10–12] ZnO presents three well studied crystal faces: a non-polar face, a Zn-polar face, and an O-polar face, labeled in Figure 1.1. A second non-polar face has also been studied, but not as extensively.[11] Both the Zn- and O- polar faces could be terminated by either a Zn or O atom. The Zn- and O- polar faces have opposite dipoles at their surfaces while the

2 Figure 1.1: Hexagonal wurtzite structure of ZnO. Gray atoms represent Zn and yellow atoms are oxygen. The three surfaces of ZnO - Zn-polar, O-polar, and non-polar are identified. The two polar faces can both either be Zn or O terminated. In this figure, the Zn-polar face is Zn terminated, while the O-polar face is also Zn terminated. non-polar face has no dipole. The surface structures in Figure 1.1 are hypothetical and are informed by extrapolation from the bulk geometry. In reality, the non-polar surface fits the bulk geometry, the Zn-terminated surface fits the bulk geometry but with apparent defects, and very clean oxygen terminated surfaces show surface reconstruction.[11] However, hyrdoxyl contaminated oxygen terminated surfaces (from water or hydrogen) do not show reconstruction, and in fact clean surfaces will convert back to the un-reconstructed surface upon contamination.[11]

The surface of ZnO is very reactive and will react with water and CO2 to form adsorbed monolayers and multilayers.[11] For this reason it is commonly used as a heterogenous catalyst, especially for methanol synthesis.[11] Water displays a rich set of adsorption structures on ZnO. As mentioned before, the oxygen terminated polar face is fully hydroxylated under small amounts of contamination in UHV.[11] The non-polar face in Figure 1.1 displays a par- tially dissociated monolayer of water at the surface in which dissociated and non-dissociated

3 molecules arrange in an alternating superstructure on the surface.[13] On nanocrystals of ZnO, IR spectroscopy showed the presence of the above mentioned adsorption structures along with other structures due to defects.[14] CO2 can react with a hydroxylated ZnO surface to form adsorbed hydrogen carbonate.[15] It is important to note here that single crystal and solution processed ZnO appear to have diﬀerent surface chemistries (regardless of crystal face), as evidenced by their diﬀering response to surface preparation techniques used in device processing.[16]

Alloying metal oxides can produce materials with tunable properties. The alloy Zn1-xMgxO is featured in Chps. 4 and5 of this thesis. Substitution of Mg for Zn in the wurtzite structure causes the bandgap to increase linearly with Mg content and the resistivity to increase exponentially, though the n-type conductivity is maintained.[17, 18] Other properties can be altered as well. For instance the resistance to acid dissolution (see Chp. 4). Other alloys of ZnO have been explored as dilute ferromagnetic oxides, where Zn is exchanged for impurity atoms intended to induce ferromagnetism, such as Zn1-xNixO, Zn1-xMnxO, Zn1-xCoxO, and

Zn1-xFexO.[19]

The films of ZnO and Zn1-xMgxO used in this thesis are produced by a sol-gel process (see Section 2.1). This process produces a very thin film of thickness 25-50 nm, as measured by profilometry. Atomic force microscopy height images of the surface of these ZnO and

Zn1-xMgxO films are shown in Figure 1.2 below. Analysis of our AFM images show that the films are composed of grains that are approximately 40 nm in size. X-ray diffraction spectra (given in Chp. 4) show that the crystal structure of these grains is the standard wurtzite structure for films up to 30% Mg. UV-Vis spectra of these films are shown in Figure 1.3 and bandgaps calculated from these spectra are given in Table 1.1.

1.2 Organic and Fullerene Semiconductors

Organic semiconductors are synthesized through the techniques of organic chemistry. They include polymers, small molecules, and dyes. Some examples are shown in Figure 1.4. There is also a class of semiconductors formed from fullerene derivatives in which fullerenes

4 Figure 1.2: Atomic force microscopy height images of Zn1-xMgxO ﬁlms produced by a sol gel process [18] for x = 0, 0.1, 0.2, 0.3. Images for x > 0 taken by Erich Meinig.

Table 1.1: Bandgaps of Zn1-xMgxO ﬁlms calculated from UV-Vis spectra. Mg Content (mol. %) Bandgap (eV) 0 3.24 5 3.33 10 3.43 15 3.52 20 3.62 30 3.77

5 Figure 1.3: UV-Vis absorption spectra of the thin ﬁlms of Zn1-xMgxO studied in this thesis. It is apparent that the bandgap increases with Mg content.

(C60,C70,C80, etc.) are modified with an attached functional group that changes the fullerene’s electronic properties or makes it soluble.[20] One of the most commonly used is phenyl-C61-butyric acid methyl ester (PCBM) shown in Figure 1.4. For the sake of simplicity, the term ’organic semiconductor’ includes fullerenes in this thesis. The physics of carrier transport and generation in these materials is highly unusual compared to inorganic semiconductors and is extremely important to the design of photovoltaic devices employing organic semiconductor materials. There are several properties that make these materials unusual: their low dielectric constants, their weak intermolecular interactions, and their very high absorption coefficients. The very high absorption coefficients of organic materials compensates for their non-ideal transport properties (discussed below).[25] While many inorganic semiconductors require micrometer-thick films to achieve complete light absorption, organic materials can generally achieve ∼95% light absorption above their absorption onset in films < 300 nm thick.[25] Thus devices can be extremely thin, making it possible to extract carriers even without

6 Figure 1.4: Examples of common electro-active organic polymers and small molecules. P3HT = poly-3-hexylthiophene. F8BT = poly(9,9-dioctylﬂuorene-alt-benzothiadiazole). CuPc = copper phthalocyanine. PCBM = phenyl-C61-butyric acid methyl ester. P3HT and F8BT images are from Ref. 21. N3 Dye is from Ref. 22. PCBM is from Ref. 23. CuPc is from Wikipedia (Ref. 24).

great transport properties. (Dye sensitized solar cells are usually much thicker because of their unique architecture, discussed in Section 1.4) However, one of the major challenges of organic semiconductors has been the design of materials with a low energy absorption onset (low bandgap in inorganic semiconductor parlance).[25] The initial materials developed all had high energy absorption onsets (≥2 eV) which is much larger than the ideal threshold energy of ∼1.3eV for terrestrial solar cells. The design of materials with a lower absorption onset is a subject of ongoing research. The relative dielectric constant (ε) of a typical organic semiconductor is ε = 4 while that of an inorganic semiconductor is ε > 10.[26] ε is inversely proportional to the strength of the electric ﬁeld within the material. Thus, the coulomb attraction between an electron and its corresponding oppositely charged hole is reduced by a factor of ε in a material. If the electron-hole pair is approximated as a hydrogenic system, the average orbital radius of the electron is increased by a factor of ε. The screening and orbital radius increase deﬁned by

7 ε scale the electron-hole binding energy by 1/ε2. Because of this, the binding energy of a photo-generated electron-hole pair in an organic semiconductor will be increased around an order of magnitude compared to inorganic semiconductors. The weak interactions between molecules in an organic semiconductor prevent them from forming strongly bound crystalline structures as are found in inorganic semiconductors. This prevents the material from forming a band structure and carriers are localized to regions on the order of one molecule in a small molecule system or several monomers of a polymer chain, instead of being delocalized as in inorganic semiconductors. The localization of carriers and the strong electron-hole interaction means that excitations in the material at ambient conditions generate bound electron-hole pairs known as excitons, with binding energies much greater than 0.1 eV, and, hence, much greater than the average thermal energy of around kT = 0.026 eV at 300 K (k – Boltzmann constant, T – temperature). This is in contrast to inorganic semiconductors which have exciton binding energies of ∼ kT or less, so that the average thermal energy is enough to split the exciton and generate free carriers. Exciton formation has important implications for organic solar cells because the excitons generated by light absorption must first be split before any current can be generated. This is typically accomplished by introducing a heterointerface, as discussed below, although excitons can also split in other ways such as at a defect in the organic. The localization of carriers also prevents band-like transport in organic semiconductors. Instead, both excitons and free carriers travel through the material via a hopping transport mechanism.[27, 28] Excitons, which are neutrally charged, can only be transported through diffusion. The diffusion length, or how far an exciton can diffuse on average before recombining, has an upper limit of 10-20 nm in most organic semiconductors used in photovoltaics. This has a significant impact on photovoltaic device design, discussed in Section 1.4. Free carriers take the form of polarons, or charges coupled to phonon modes.[29] The basic idea is that the presence of the free charge causes a relaxation of the surrounding atoms. This relaxation can be significant in organic semiconductors because these materials have

8 relatively labile bonds that distort easily, as compared to inorganic semiconductors. This distortion of molecular structures follows the charge as it moves through the material. Free carrier transport is also dramatically influenced by the degree of order of the molecules in the film and the overall film morphology. Grain boundaries, stacking and packing geometry, and crystallinity can all affect transport dramatically.[30–32] A distinction must be made about electronic structure in organic semiconductors compared to inorganic semiconductors. Because of weak intermolecular interactions, there are no valence band or conduction band states. Instead, transitions occur between localized molecular orbitals defined by the structure of the molecule and its interaction with neigh- boring molecules. The minimum energy electronic transition is therefore between the highest energy occupied molecular orbital (HOMO) and the lowest energy unoccupied molecular orbital (LUMO). These are analogous to, but distinctly different from, the valence band and conduction band of crystalline inorganic semiconductors. In a well ordered organic material, interactions between molecules can give rise to multi-molecule excitations that are lower in energy than the HOMO-LUMO gap of the isolated molecules. This can be observed, for instance, in well ordered P3HT, and can induce significant changes in the absorption spectrum, in addition to the impact on charge transport mentioned above.[33] Organic semiconductors often form the active layer of the photovoltaic devices they are used in, meaning they are often absorbing light and separating charge. Because of the importance of excitons in these devices, they are often referred to as excitonic devices. In these devices, it is necessary to split the exciton before the rest of the photovoltaic process can occur. The mechanism that typically accomplishes this in these devices is an ultra-fast charge transfer process at the interface between two organic semiconductors or an organic semiconductor and inorganic semiconductor. The material that gives up an electron is referred to as a donor and the material that receives it is referred to as an acceptor. This charge transfer occurs at the interface between two carefully matched materials. A minimum condition for charge transfer is that the free energy gained in the charge transfer process must

9 be high enough to overcome the exciton binding energy. There must be some energy offset at the interface of these materials to allow this to occur, for instance a difference in the LUMO energy of the two materials. In standard semiconductor physics, this is referred to as a type II heterojunction.[34] The discovery of ultra-fast electron transfer from a semiconducting polymer or small molecule (the donor) to a fullerene derivative (the acceptor) is exploited in nearly all organic photovoltaic devices to date.[35] The charge transfer rate from polymer to fullerene has been measured to be less than 50 fs, far faster than any recombination process.[36] Similarly, charge transfer from a dye molecule attached to a metal oxide scaffold in a dye sensitized solar cell shows a fast charge transfer component occurring in less than 100 fs, again several orders of magnitude faster than any recombination processes, and a slower picosecond process.[37]

1.3 General Photovoltaic Device Physics

Excitonic solar cells have a number of specialized requirements that differentiate them from inorganic solar cells. All of these design requirements derive from the property of organic semiconductors, discussed in Section 1.2, that light absorption generates excitons instead of free carriers. This requires a specialized architecture for generating free carriers, but upon free carrier generation the same general theory used for inorganic solar cells can be applied to excitonic devices. For this reason a general discussion of solar cell device physics is provided first, followed by discussion of the new ideas required to understand excitonic devices in Section 1.4. Solar cells are essentially layered structures composed of conducting and semiconducting materials. In isolation, each of these materials has a thermal equilibrium distribution of electrons and holes. This thermal distribution can be described by an absolute chemical potential energy referred to as the Fermi energy, Ef , in semiconductor physics. Further explanation about how the Fermi energy describes semiconductors in equilibrium can be found in the introductory text by Neamen.[34] The chemical potential energy provides a measure of the potential for charge carriers to diffuse. So when the isolated materials are

10 brought together in a solar cell, charge will diﬀuse between them according to their diﬀerences

in Ef until a new equilibrium is established and Ef is constant across the whole device. However, the redistribution of charge sets up electric ﬁelds in the device that can be described by the gradient of the familiar electric potential, Φ. The equilibrium net electric potential

across the device is referred to as the built-in potential, Vbi. This process is illustrated in Figure 1.5 below for a standard pn-junction and for a metal-insulator-metal junction, which is commonly used as a qualitative model of organic photovoltaic devices.[38]

Figure 1.5: Process of Fermi level equilibration in junctions of electronic materials. (a) Stan- dard pn-junction equilibrium. Before contact (i), the p and n-type materials have differing Fermi levels Ef,1 and Ef,2. After the junction is formed (ii), the Fermi levels equilibrate by charge transfer and band bending to achieve a constant Fermi level, Ef . During this process a built-in electric field develops, with potential drop Vbi. (b) In the metal-insulator-metal model of organic solar cells, the active layer is approximated as insulating. Before contact (i), the metal contacts have differing Fermi levels. After contact (ii), the Fermi levels of the metals equilibrate through metal-to-metal charge transfer. This charge transfer generates a built-in field and potential across the insulating active layer.

11 The photovoltaic eﬀect is the generation of a current in a device when light is incident on it, without the need to apply an external voltage, also implying the generation of a photovoltage. In order for current to be produced, there must be an asymmetry in the device that drives the electrons and holes generated by the light in opposite directions. In solar cells this can be accomplished in many ways, for instance the donor/acceptor structure in excitonic devices, discussed further below, or the existence of a built-in potential Vbi, as in a conventional pn-junction solar cell. When light is incident on the device, thermal equilibrium is disrupted, and a non-thermal distribution of carriers results. The thermal distribution of carriers described by Ef is no longer valid. Excess electrons and holes are generated in the conduction and valence states, respectively. Luckily, these carriers generally equilibrate to a thermal-like distribution in their respective states at a rate much faster than they recombine with the opposite carrier.[39] This makes it possible to deﬁne a quasi-Fermi (quasi-thermal) energy for both electrons and holes separately [26]: n Ef,n = Ec + kT ln( ) (1.1) Nc p Ef,p = Ev + kT ln( ) (1.2) Nv

Ef,nand Ef,p are the quasi-Fermi energies for electrons and holes respectively. Ec and Ev are the energies of the conduction states (c), occupied by electrons, and valence states (v), occupied by holes, of the materials in the device. n and p are the electron and hole concentrations, respectively, in the conduction and valence states. Nc and Nv are the density of states of the lowest energy conducting states and highest energy valence states. These equations are intended to be generally descriptive of either organic or inorganic semiconductors. However, the interpretations of Ec, Ev, Nc, and Nv are complicated by the presence of disorder in a material which spreads the density of states out so that these quantities are less clearly deﬁned. This disorder has signiﬁcant implications for excitonic solar cell performance.[40] Each quasi-Fermi energy is actually composed of the familiar potential energies already discussed above. Taking Ef,n as an example, Ec is a function of position and

12 follows the variations in the electrostatic potential energy eΦ through what is commonly referred to as band-bending.[34] The term kT ln(n/Nc) captures the chemical potential because it is a function of electron concentration, n, which varies with position.

The quasi-Fermi levels Ef,n and Ef,p really define the behavior of the device. They shift when an external load is applied that alters the electrostatic potential. They shift as the intensity of light changes or if the rate of electron and hole recombination goes up. The gradients of Ef,n and Ef,p define the forces on the electrons and holes, respectively, and therefore define the current in the device [26, 34]:

J = nµn∇Ef,n + pµp∇Ef,p (1.3) or equivalently,

J = −e (nµn + pµp)∇Φ + ekT (µn∇n − µp∇p) (1.4)

J is the current density, e is the fundamental unit of charge, p and n are the hole and electron volume densities, respectively, k is Boltzmann’s constant, T is temperature, µp and

µn are the hole and electron mobilities, respectively, Φ is the electric potential, and ∇ is the gradient operator. These equations can be derived from the relaxation time approximation of the Boltzmann transport equation (or more simply the Drude model) and the Einstein relation, which has been employed here to give the electron and hole diffusion coefficients as kT µn,p.[34, 39] The values of Ef,n and Ef,p at the contacts to the device define the amount of usable energy available from a pair of carriers, and therefore the voltage (V ) of the device [26]:

− + eV = Ef,n (e contact) − Ef,p (h contact) (1.5)

These ideas make it possible to understand the meaning of the performance metrics reported for solar cells. These parameters are illustrated in Figure 1.6, where an example current density-voltage (J-V) curve is shown. The current density J is the current per unit area flowing through the device, defined by Equations 1.3 and 1.4. V is the voltage across the load on the device. The voltage generated in the solar cell, defined by Equation 1.5, will

13 match the voltage drop across the load, V , by Kirchoﬀ’s law.

Figure 1.6: Example of a current density (J) - voltage (V) curve for a solar cell. The important performance metrics are shown.

The shape of the J-V curve is due to the fact that good solar cells are rectifiers, only allowing current to flow in one direction. This is a consequence of the asymmetric, charge separating character of the device that allows photovoltaic action, discussed above. When the load is zero, the device has been shorted. Though V = 0, gradients in the quasi-Fermi energies still exist that drive the carriers, resulting in a photo-generated current. This is known as the short circuit current, Jsc. At the other extreme, when J = 0, the circuit is effectively open and the voltage under this condition is referred to as the open circuit voltage

(V oc). Because the current switches direction at this point and power cannot be extracted from positive currents, eVoc represents the highest energy per electron that can be collected.

V oc is highly sensitive to a number of very important processes occurring in the device. Whereas J is sensitive to gradients in the quasi-Fermi energies, V is sensitive to their actual values. Combining Equations 1.1, 1.2, and 1.5 and assuming that J = 0, we ﬁnd

14 n p eVoc = Ec + kT ln( ) − Ev − kT ln( ) (1.6) Nc e−contact Nv h+contact The brackets indicate which contact each of the terms should be evaluated at. The ﬁrst thing

to notice is that Voc depends on the diﬀerence in the energy of the charge transporting states

− + of each of the contacts, Ec(e contact)−Ev(h contact). This makes it clear that the electronic structure of the contacts to a solar cell, and its interaction with the active layer, can be quite important. This will be the focus of Chp. 3. Throughout the device, the conducting state

energies, Ec and Ev, deﬁne a ’baseline’ for the quasi-Fermi level and their diﬀerence should

be as large as possible in order to maximize Voc. The carrier concentrations n and p also play a signiﬁcant role in determining Voc. n and p are inﬂuenced by carrier generation processes (light absorption and exciton splitting in excitonic solar cells) and carrier recombination processes and thus Voc is a sensitive measure of both. The higher n and p, the better, and eliminating recombination processes that reduce the carrier concentration is an important strategy for maximizing current and voltage, a somewhat counter-intuitive result, but a very

important one.[41] The carrier concentration terms in Equation 1.6 also suggest that a Voc can be generated even if there is no difference in the energy of the contacts, purely through a diffusion force.[26, 42] Theoretical calculations also suggest that a diffusional force can counteract a loss of energy due to non-optimal contacts in bilayer organic solar cells.[43]

Along the J-V curve, between V = 0 and Voc, there is a maximum power point (MPP)

with PMPP = JMPP VMPP (Figure 1.6), where P is the power delivered by the solar cell. Note that this is a power density (power per unit area). The MPP can be used to calculate the eﬃciency of the device (η) from: P η = MPP (1.7) ILight

where ILight is the power available from the incident light. Two rectangles are shown in

Figure 1.6, one representing the actual maximum power generated, with area PMPP , and

one representing the ideal maximum power, with area JscVoc. The ratio of these is called the ﬁll factor (FF ):

15 P FF = MPP (1.8) JscVoc The FF is a quantification of how closely a device approaches ideal power generation, and a quantification of how current decreases as the load (bias) on the device increases. The FF characterizes the recombination processes that are manifested as the load on the device is increased. The loss processes that contribute to the FF , and to device performance in general, are a topic of current research in excitonic solar cells.[44, 45] In order to understand the discussion of device performance in Chp. 3 and in the scientific literature in general, it’s important to understand the equivalent circuit model of a solar cell shown in Figure 1.7. The device is modeled in terms of three elements: a rectifying junction

(or diode, D), a resistance in parallel, the shunt resistance (RSH ), and a resistance in series

(RS). The diode is modeled by the ideal diode equation, sometimes with modiﬁcations for non-ideality.[34] The photocurrent is modeled as a voltage-independent current source IPH in parallel with the diode. The circuit is completed by the load, which has a voltage drop V across it (implying a bias of V across the solar cell). The major advantage of this model is that it provides a simple way of incorporating loss processes into a solar cell model. All the device performance metrics discussed above (Figure 1.6) can be derived from a standard circuit analysis. The loss processes are incorporated through RS, which represents the Ohmic losses accrued from resistance within the device, and RSH , which represents the carrier concentration losses due to recombination and any shorts that might exist in the device. These losses can be approximated from experimental J-V curves and are often reported in the literature. In Chp. 3, RS was calculated from the slope of the J-V curve at V = 1V and

RSH was calculated from the slope at V = 0V. While this equivalent circuit model is simple and easy to use for achieving a general understanding of experimental data, more complex modeling is required to accurately describe the J-V curves that are encountered in real solar cells.

16 Figure 1.7: Equivalent circuit model of a solar cell. The device is composed of a voltage- independent current source, IPH , due to the incident radiation, a rectifying junction (D), a shunt resistance (RSH ) representing shorting and recombination processes, and a series resistance (RS) representing Ohmic losses in the device. The device is attached to a load with voltage drop V across it.

1.4 Excitonic Solar Cell Device Physics

The preceding concerns the behavior of free carriers in solar cells. However, generation of free carriers upon light absorption is not guaranteed. Excitonic solar cells employ materials in which light absorption generates excitons instead of free carriers. This requires the architecture of these devices to deviate from that of conventional solar cells. As discussed in Section 1.2, it is necessary to build the active layer from two diﬀerent materials: a donor and acceptor. The materials are chosen in order to allow excitons to eﬃciently dissociate into free carriers at the interface of these two materials. There are four fundamental steps in the operation of an excitonic solar cell (Figure 1.8)[46]:

1. Generation: Excitons are generated by light absorption.

2. Diﬀusion: Excitons diﬀuse to an interface.

3. Exciton Dissociation: the exciton is split by charge transfer across the interface (as discussed in Section 1.2), and the charges are free to move.

4. Charge Collection: charges are driven by concentration gradients and built-in electric ﬁelds to the contacts where they are collected and used to power a load.

17 Figure 1.8: The operational steps of excitonic solar cells. Figure based on Ref. 46. The electron and hole reside below the energy of the HOMO-LUMO optical gap in steps 1-3 due to their exciton binding energy.

In order to maximize exciton dissociation and light absorption simultaneously an architecture unique to excitonic photovoltaics has been developed. The donor and acceptor must be intimately blended at the nanoscale so that excitons can diffuse to an interface without recombining. The exciton diffusion length of ∼10 nm defines the length scale of this blending.[46, 47] This design allows the film to be optically thick while still allowing for efficient exciton dissociation. Furthermore, a percolation pathway must exist in the donor and acceptor material phases that allows free carriers to leave the active layer efficiently. Examples of the specialized architecture required are shown in Figure 1.9. In Figure 1.9(a), the polymer P3HT is combined with PCBM in a ’bulk heterojunction’ in which domains of P3HT and PCBM form a large interface area for exciton dissociation. Figure 1.9(b) shows a hybrid nanowire ZnO/P3HT photovoltaic device. The high interface area created by the P3HT/nanowire interface is intended to efficiently split excitons. However, careful investigation of these devices suggests that exciton dissociation at the metal oxide/organic interface

18 is very inefficient, and that most of the free carriers generated are due to exciton splitting at defects in the P3HT.[48] Figure 1.9(c) shows a dye sensitized solar cell (DSSC). In these devices a light-absorbing dye is chemically bound to a high surface area metal oxide such as a porous matrix or nanowires.[37] An exciton generated in the dye is efficiently split by charge transfer to the metal oxide. This design eliminates step two, exciton diffusion, in the operation of an excitonic solar cell. High light absorption is achieved because of the extremely high surface area of the porous metal oxide matrix.

Figure 1.9: Examples of excitonic photovoltaic device architectures. (a) Inverted bulk heterojunction solar cell. (b) Metal oxide/organic nanowire solar cell. (c) Nanowire dye sensitized solar cell (DSSC). The ﬁgures are meant to be illustrative and are not drawn to scale.

After the exciton has been split, free carriers (polarons) are generated. From Section 1.3, the two processes that drive carriers to the contacts during charge collection are drift and diffusion. Diffusion plays an important role in excitonic solar cells because exciton dissociation creates a high concentration of each carrier type in physically separated regions.[26] The concentration gradient developed at the exciton-splitting interface causes carriers to diffuse away from it toward the contacts, aiding in charge separation and transport. Thus, in these devices the energy offset between donor and acceptor helps to drive efficient charge separation and transport, in addition to separating excitons. The role of drift in charge transport differs between different device designs. In bulk heterojunction and metal oxide/organic

19 heterojunction solar cells, fields contributing to drift can only really originate in the Fermi level difference between the contacts. The reason for this is that, while electric fields may exist locally along the donor-acceptor interface, the complex, interdigitated architecture of these devices, along with low carrier concentrations in a very thin active layer, rules out the possibility of a built-in electric field being established by the active layer itself. However, differences in the Fermi level of the contact materials can establish a built-in potential Vbi that extends across the active layer, contributing to drift. In DSSCs the situation is substantially different. The liquid electrolyte permeating the active layer will completely cancel any external electric fields and in this case carrier transport is governed entirely by diffusion.[42] There are a number of unsolved questions about the physics of excitonic devices, especially in the newer organic photovoltaic devices. As mentioned above, recombination in these devices is an area of active research. The origin of the open circuit voltage is still an open question as well. A number of studies have established that the energy difference between the HOMO of the donor and the LUMO of the acceptor in optimized devices has a 1 V/eV influence on Voc.[18, 49, 50] However, the influence of the contacts on Voc in these devices is not well understood. The discussion in Chp. 3 reviews this question extensively. Another area of recent interest in organic photovoltaics is contact selectivity.[51] In order to prevent recombination at the contacts and retain carrier density in the active layer, it is best if each contact selectively passes or transports only one carrier type. This is of general interest to all of photovoltaics, however the impacts of it on organic solar cells are just beginning to be studied.[52] These initial results suggest that charge selective contacts create an enhancement in Voc, likely through maintaining higher carrier densities near the contacts. Metal oxide contacts such as n-type ZnO or p-type NiO are expected to be selective because the materials are intrinsically doped to transport one carrier type, and large energy barriers exist for the extraction of the opposite carrier (see Section 1.5). ITO, on the other hand, is not expected to be selective because of its metallic character.

20 1.5 Metal Oxide/Organic Interfaces

Metal oxides are employed in organic electronic devices as contacts to the organic active layer of the device (as in a bulk heterojunction solar cell) or as a component of the active layer itself (as in hybrid photovoltaic devices or dye sensitized solar cells).[37, 51, 53] The performance of the interface between these two layers is critical to overall device performance, and understanding and optimizing these interfaces is an important research direction in excitonic photovoltaics and electronics in general. There are three major areas of interest:

1. Morphology of the organic layer at the interface. Does the organic phase wet the metal oxide, and how well does the organic phase order near the interface with the metal oxide?

2. Energy level alignment: how do the energy levels of each component shift after the interface is formed and how do the orbitals/bands of each component interact with each other? Are there energy barriers formed that could inhibit charge transfer?

3. Chemical reactivity: do chemical reactions occur between the organic and the metal oxide, and how does this aﬀect interface performance?

Examples of these ideas are illustrated in Figure 1.10. Figure 1.10(a) shows how the metal oxide/organic interface can impact the morphology of the organic, potentially leading to poor charge transport. This has been observed in the case of ZnO/P3HT interfaces, where P3HT near the interface shows a blue shift in its absorbance spectrum due to reduced interchain interactions.[54, 55] In Figure 1.10(b), holes are unable to transfer from the organic to ZnO due to the large diﬀerence in energy between the organic HOMO energy and ZnO’s valence band energy. On the other hand, electrons can transfer easily. In this case, the electronic structure makes the ZnO electron-selective.[51] Figure 1.10(c) shows how dye molecules can be chemically attached to a metal oxide through a carboxylate bond. In this case, a chemical bond has been purposefully introduced to anchor the dye close to the surface, allowing charge

21 transfer and photovoltaic action to occur. In other cases, unexpected chemical reactions between the interface components may be detrimental.

Figure 1.10: Examples illustrating metal oxide/organic semiconductor interface properties important to organic photovoltaic device performance. (a) The organic morphology may become disordered at the interface, preventing eﬃcient charge and exciton transport near the interface. (b) Interface electronic structure can give charge selective contacts. (c) Chemical bonding of dye to metal oxide makes eﬃcient charge transfer possible.

The electronic structure of the interface is a major topic of Chp. 3 and will be the focus of the rest of this section. Figure 1.11 shows the electronic structure of an isolated n-type metal oxide (n-MO), an organic, and two potential interfaces in more detail. Energy levels in materials are often referenced to what is referred to as the vacuum level or vacuum energy, shown as the high energy solid black line in Figure 1.11. This is the energy of a free electron that has just barely escaped the potential of the material, and resides just outside the material.[56] Theoretically a vacuum level can be defined within a material, but it is not experimentally accessible. It is a useful reference, however, because it tracks changes in the electrostatic potential which a free electron will still experience. The energy required to extract an electron from the Fermi energy at the surface of the material (Ef,surface) to this vacuum level is called the work function (φ in Figure 1.11) and is a useful quantity for comparing materials because differences in the work function indicate differences in the chemical potential of the carriers within each material (Section 1.3), and because it can be measured experimentally. Also important are the differences in energy between conducting

22 states on either side of the interface. In the case of Figure 1.11, this is the diﬀerence between

the LUMO level of the organic and the conduction band of the metal oxide (ELC ). When comparing isolated materials (Figure 1.11(a)), this gives an estimate of what the energy barriers to charge transfer will be.

Figure 1.11: Electronic structure of an n-type metal oxide and an organic in isolation (a), after an interface is formed under assumption of vacuum level alignment (b), and after an interface is formed in which there is a dipole-forming charge transfer or chemical interactions between the two components (c). n-MO means n-type metal oxide, φ is the work function of the interface, ∆φ is a dipole-induced work function change, CB is conduction band of the metal oxide, Ef is the Fermi energy, E is energy, ELC is the energy diﬀerence between the LUMO of the organic and the conduction band of the metal oxide.

Upon actual formation of an interface, non-idealities are expected. It is useful to ﬁrst discuss the ideal situation. The ideal energy level alignment, which is often depicted in the excitonic photovoltaics literature, is shown in Figure 1.11(b), and is referred to as vacuum level alignment. In this situation, the vacuum level remains constant as in the case of isolated ﬁlms, and the work function of the interface is the work function of the substrate. Vacuum level alignment has two major assumptions: (1) the materials are non-interacting, i.e. there are no chemical reactions between them and (2) the organic behaves like an insulator, and charge transfer between the organic and substrate is inhibited.[57] These approximations are surprisingly good in metal and metal oxide/organic interfaces where the metal oxide work function is in the HOMO-LUMO gap of the organic.[57–59] Braun et al. suggest charge

23 transfer between the organic and substrate is inhibited by a charge-blocking contamination layer due to sample preparation in an ambient environment.[57] However, Greiner et al. have prepared metal oxide/organic interfaces under ultra high vacuum conditions by evaporating an organic layer on top of a metal oxide layer oxidized within the chamber.[59] They still observe vacuum level alignment and propose an alternative explanation in which the molecular nature of the organic prevents charge transfer unless it is thermodynamically stable for a molecule to be charged.[59] Explaining the vacuum level alignment observed in these experiments is still a matter of fundamental research. Violation of (1) and (2) invariably induces an interface dipole that can significantly alter the energy level alignment at the interface compared to what would be predicted in the vacuum alignment case. This situation is depicted in Figure 1.11(c). The dipole can be approximated as two infinite sheets of charge. A simple application of Gauss’ law shows that this arrangement has an electric field of zero on either side of the dipole and a constant electric field between the charges, leading to a linear step in electric potential energy across the dipole (∆φ). Whether the step raises or lowers the energy depends on the direction of the dipole. This dipolar energy step is manifested as a shift in the vacuum level, as even free electrons will experience this dipole. Electrons leaving the material must traverse this dipole and the work function of the film changes by an amount ∆φ which is the potential energy shift due to the dipole. Consequently ELC is shifted by the same amount. One of the most dramatic observations of charge transfer and subsequent dipole formation at the interface is referred to as Fermi level pinning. Fermi level pinning occurs when a high density of states are present in one material at an interface. In this case, charge transfer occurs to or from these states until the Fermi energy reaches the energy of these states.[60] In metal oxide (and metal)/organic interfaces, this is observed when the Fermi level of the metal oxide approaches either the LUMO or HOMO level of the organic, where a high density of states is available.[57–59] The Fermi level of the interface becomes pinned to the HOMO or LUMO level and doesn’t change no matter how high or low, respectively, the work function

24 is. When considered with the discussion above, the overall picture of metal oxide/organic interface alignment is that vacuum level alignment is observed when the metal oxide Fermi level is in the HOMO-LUMO gap, while strong Fermi level pinning is observed near the LUMO and HOMO. Excitonic devices are quite sensitive to energy level alignment across metal oxide/organic interfaces. In DSSCs, the LUMO of the dye must be higher in energy than the metal oxide conduction band in order for charge transfer to occur. In hybrid devices, where the active layer is made up of a metal oxide/organic interface, the energy level alignment has been shown to have a very strong inﬂuence on Voc.[18] In an experiment by Olson et al., where the metal oxide conduction band energy was varied using Zn1-xMgxO alloys, the Voc showed a

1 V/eV correlation with the conduction band edge.[18] Jsc also showed a strong correlation with the conduction band edge as well. This is in agreement with the results observed for all-organic active layers (Section 1.4). In organic bulk heterojunction and planar devices, where metal oxides are used as charge selective contacts, the situation is substantially more complicated. The relationship between energy level alignment between the contact and active layer and device properties is not well understood. Chp. 3 focuses a lot of attention on this issue. While metal oxides already play an important role in excitonic photovoltaics, the catalog of metal oxides available isn’t necessarily enough to satisfy the requirements of every device, and there are issues that can’t necessarily be solved by changing the metal oxide bulk properties. To that end, further tunability of the interface is desired. One of the most promising techniques, surface modiﬁcation using molecular monolayers, is the main strategy discussed throughout this thesis.

1.6 Monolayer Modiﬁcation of Metal Oxide Surfaces and Interfaces

An important strategy for trying to optimize metal oxide/organic interfaces in optoelectronic devices is surface modiﬁcation of the metal oxide. The idea is to alter the physical properties of the metal oxide surface in order to optimize an interface. There are a couple

25 ways to approach this. The first is to chemically alter the surface through cleaning processes such as UV-ozone cleaning or oxygen plasma cleaning. There have been a number of studies showing the impacts of these treatments on devices.[61–64] However, there are reports that these techniques are not stable and don’t permanently alter the surface.[61, 65] A more stable technique, that presents substantially more tunability, is the attachment of molecular monolayers to the surface of the metal oxide. The monolayer is made up of molecules that are anchored to the surface through an attachment group such as a phosphonic or carboxylic acid, thiol, or triethoxysilane.[54, 66– 69] These chemistries are chosen to make the attachment process surface selective to prevent multilayer formation. The attachment group anchors a functional group to the surface. The exposed functional group is chosen to have attributes that achieve the desired modification of the surface properties. The uses of molecular monolayers are quite varied and show that they are a versatile technique. In Section 1.5 the disordering of semiconducting polymers at the ZnO interface was described. Attachment of alkyl monolayers via either a thiol or triethoxysilane chemistry can prevent the adverse interaction between metal oxide and polymer that disrupts polymer ordering.[54, 55] In the case of the alkanethiol, this led to improvement of hybrid ZnO/P3HT solar cells.[54] It is also possible to introduce dipoles at the interface, as described in Section 1.5, by attaching molecules with built-in dipoles. This can improve energy level alignment. This is the subject of Chp. 3 and is discussed in-depth there. Monolayer modification plays a significant role in DSSCs as well. It’s clear that dye attachment itself is a form of monolayer modification. Recently, however, co-adsorption of a dye and a molecular modifier has contributed to improved device performance. The co-adsorbant can have two effects. One effect is to prevent aggregation of dye in which this aggregation is common.[37] The other effect is to reduce recombination.[2] Recombination is not a big issue when the iodide/triodide electrolyte solution is used, but for solid-state hole conductors and other electrolyte solutions it is a significant problem. The use of molecular modifiers to separate the electrolyte and metal

26 oxide or to induce dipoles at the surface can significantly slow the rate of recombination.[2] The attachment schemes used to modify metal oxides are usually based on acid-base reactions between an acidic modifier and basic surface. The attachment process begins with protonation of the surface, followed by complexation of the metal atoms by the anionic modifier. In ZnO this can lead to significant dissolution (see Section 1.7) and for this reason Chp. 3 explores the use of triethoxysilane modifiers that bond via a covalent Si-O bond, while Chps. 4 and 5 consider an alternative to ZnO.

1.7 Acid Dissolution of Metal Oxide Semiconductors

Many of the attachment schemes used to modify metal oxide surfaces involve the use of acidic attachment groups, for instance phosphonic and carboxylic acids. ZnO is susceptible to etching even by weak acids and therefore some destruction of the ZnO surface is inevitable with organic acid modiﬁers. This has been observed for a number of diﬀerent attachment species.[8, 9, 70] This problem is a central theme of this thesis, and a general understanding of how metal oxides are dissolved by acids is quite important. The metal oxide surface is comprised of metal and oxygen atoms bonded to each other (Me-O-Me or oxo groups). Dissolution of metal oxides by acids is accomplished by three general steps[71, 72]:

1. Protonation (attachment of protons) of surface oxo groups

2. Rupture of Me-O bond

3. Phase transfer of the metal from the surface to solution

This process can occur in a variety of ways depending on the metal oxide and the nature of the acidic dissolving solution, however Me-O bonds must be broken by either nucleophilic or electrophilic attack or both.[71] Electrophilic attack involves reaction of the oxygen atom and is accomplished by protons (H+) in acidic solutions. Nucleophilic attack involves reaction of the metal with an anionic species such as a ﬂuorine ion (F-) or a complexing species.[71]

27 Based on this reaction process, the surface charge due to adsorbed protons, hydroxyls, and complexing species is clearly of importance to dissolution of metal oxides by acids. The net charge will determine the rate at which the dissolution reaction can occur because it determines the surface concentration of etch species. Metal oxides accumulate surface charge depending on their particular surface chemistry, and this will influence their rate of dissolution in acid. The surface charge can be quantifed in the simplest picture of acid dissolution, in which only protons and hydroxyl groups adsorb to the surface, by the point of zero charge (pzc) of the metal oxide. The pzc is the point at which surface hydroxide and surface proton concentrations are balanced.[71] The pzc occurs at different values of pH (referred to as pHpzc) for different materials, and determines net excess of protons or hydroxyls at a particular pH for a given material. The material will have an excess of protons at pH below pHpzc and an excess of hydroxyls above. As an example, consider ZnO with pHpzc of around 9 and TiO2 with pHpzc of around 6.[70] Under acidic conditions the ZnO surface will have a large imbalance of protons vs. hydroxide, while this imbalance will be substantially reduced for TiO2. While the pzc will certainly influence etching through its determination of surface proton concentration, it is not the defining variable in the etch rate. The reactivity of the Me-O bond is also very important. This is discussed further in Section 4.3 as it pertains to the model of metal oxide etching presented there. It is important to note that the pzc assumes an ideal situation in which only protons and hydroxide ions contribute to the surface charge. However, complexation by other ions is very common, and indeed very important to acid/base surface modification schemes. Under conditions of multiple adsorbing ions, the pH at which net surface charge is zero deviates from pHpzc, and the new point of zero net charge is referred to as the isoelectric point (IEP).[71] The role of complexing species can be instrumental to metal oxide dissolution. Since this thesis centers around formation of surface complexes by molecular modifiers, the role of complexing species must be discussed. This is done in the context of semiconducting metal oxides of interest in this thesis. There are two regimes of semiconductor metal oxide

28 dissolution. In the first regime, the dissolution is dominated by the electronic structure of the surface of the metal oxide. This is common with non-adsorbing mineral acids (HCl, etc.).[71] The second regime occurs in the case of highly adsorbing dissolving species. In this case, the dissolution process is dominated by the surface complexes formed, which are electronically decoupled from the electronic structure of the solid surface.[71] This case is likely apt to describe dissolution of metal oxides by organic acid surface modifiers since they are strongly adsorbing. In this case the rate of dissolution is limited by the rate of transfer of surface complexes to solution.[72] It is expected that this is the rate-limiting step of dissolution by our surface modifiers, and any changes seen in the etch rate are related to this transfer rate. The transfer of complexes from the surface to solution also implies the production of metal-organic complexes of the metal and the adsorber anion. These can precipitate onto the surface, and can be a significant problem, especially in ZnO-based DSSCs.[9, 70]

1.8 Thesis Organization

This thesis is about overcoming the issue of dissolution associated with organic acid modification of ZnO. Two major strategies will be discussed: the use of covalently bonded modifiers that don’t etch ZnO and the substitution of Mg for Zn to create Zn1-xMgxO alloys which display higher resistance to etching by acids. Before discussing the results of the research, Chp. 2 will give some background on the methods employed to produce the materials and surfaces of interest and the techniques used to characterize them. Chp. 3 will discuss the use of molecular surface modifiers that bind through the triethoxysilane chemistry to alter the surface properties of ZnO. The triethoxysilane attachment chemistry forms a covalent bond to the surface, replacing the commonly used acid base chemistry. This chapter contains a complete device study relating to the questions arising in Sections 1.5 and 1.6 about the origin of Voc in bulk heterojunction devices. It shows definitively the usefulness of the triethoxysilane chemistry in tuning interface properties. Chp. 4 discusses surface modification and etching of Zn1-xMgxO alloys by benzoic acid, a prototypical carboxylic acid modifier. The etch resistance of Zn1-xMgxO is demonstrated,

29 and possible mechanisms are presented. The surface species created by exposure to benzoic

acid are also identiﬁed. Chp. 5 extends the study of Zn1-xMgxO to N3 dye, commonly used in DSSCs, that bonds through a carboxylic acid group. This study discusses the initial etch rate – the rate at which the ﬁrst few atomic layers are etched - providing useful information for evaluating the etching mechanisms hypothesized in Chp. 4. This chapter also explores

the idea that the improved etch resistance of Zn1-xMgxO can reduce the rate of formation of Zn-dye complexes that are produced during the sensitization process. Finally, Chp. 6 draws conclusions about the results of this thesis and suggests future research paths that are motivated by this work.

30 CHAPTER 2 SAMPLE PREPARATION AND CHARACTERIZATION TECHNIQUES

Rather than supplying recipes and procedures (which can be found in the experimental sections of each results chapter), this chapter aims to illuminate the underlying physical and chemical processes that the techniques used in this thesis are based on.

2.1 Production of Zn1-xMgxO Thin Films

The thin films described in this thesis have been produced by a sol-gel process. In a sol-gel process, molecular precursors are first suspended in a colloidal solution (the sol). Then the solution is cast into a film. A drying step then condenses the film into a much more dense gel. Finally, heating is applied in order to react the precursors to form the final compound. In the case of the films in this thesis, the drying and reaction step are preformed simultaneously.

Zn1-xMgxO films can be produced from a sol of Zn and Mg acetate.[18, 73] The basic idea is that the Zn and Mg acetate film will decompose in air. The acetate (the deprotonated form of acetic acid) decomposes into carbon dioxide and acetone, leaving behind one oxygen with which the oxide is formed. These films are considered metastable. The equilibrium phase diagram for ZnO and MgO suggests the solubility of these two compounds is only around 4%, and high temperature annealing can cause phase separation.[74] This means that the films formed in the sol-gel process are not at equilibrium and it is expected that the processing conditions impact the film quality substantially.[75] The films should be dehydrated and decomposed quickly in order to avoid phase segregation.[75] We make a note here about the chemicals used to form our particular films. We use solutions of Zn acetate and Mg acetate in 2-methoxyethanol with ethanolamine. The ethanolamine

31 is air sensitive and it is best to store it in a glove box until it is needed. Outside the glove- box, ethanolamine can last for several months before eventually changing from a clear color to yellow, indicating the formation of some other chemical compound(s). We recommend changing to new ethanolamine periodically before this color change occurs, as the yellow liquid produces inferior ﬁlms.

2.2 Triethoxysilane Modiﬁcation of ZnO

Silanes form a broad class of chemicals that can bond through a Si-O bond to a surface. The reaction requires the presence of hydroxyl (OH) groups on the surface and was ﬁrst

developed for SiO2, whose surface is hydroxylated under ambient conditions.[76] Many of these compounds are extremely reactive to water, polymerizing into siloxane networks in solution and producing multilayers on the surface.[76] However, it has been found that this process occurs much more slowly in triethoxysilanes (TES), making them suitable for forming monolayers in ambient, ’wet’ conditions.[69, 77] Because of the low reactivity, a catalyst, n- butylamine, is used to expedite the surface attachment process. The addition of the catalyst doesn’t accelerate the reaction of triethoxysilanes in solution, but only near the surface.[78] The process of attaching triethoxysilanes to ZnO has been successfully transferred from

SiO2, with only minor modiﬁcation.[55, 69, 79] The task of hydroxylating ZnO presented some diﬃculty. The solution came from previous reports that ultraviolet light can induce hydroxylation of the ZnO surface.[80] We have found that UV-ozone cleaning, in which ozone is created by ultraviolet light, can successfully hydroxylate ZnO.[69] This UV-ozone clean comes at a cost to the surface quality: UV-ozone cleaning dramatically reduces the

Voc of bilayer P3HT/ZnO photovoltaic devices.[62] UV-ozone cleaning appears to induce substantial changes in surface electronic properties, including increased work function and reduced ionization potential.[16, 62] We have suggested that these changes may carry through

TES treatment, resulting in lower Voc in devices with TES treated ZnO as well.[55, 79] The TES attachment scheme produces disordered monolayers that have sub-monolayer coverage as compared to a mechanically assembled Langmuir-Blodgett ﬁlm produced from

32 a similar molecule.[69] While the TES molecules are expected to cross-link on the surface through their remaining ethoxy arms, to form Si-O-Si bonds, no spectroscopic evidence has been found to support this. However, Si-O-Eth (Si-ethoxy) bonds have also not been detected so at this point it is not known if these molecules are cross-linked, though this is expected based on the attachment chemistry.

2.3 Carboxylic Acid Modiﬁcation of Zn1-xMgxO

As discussed in Section 1.6, acid-base reactions between acidic modiﬁers and metal oxide surfaces have been very important to tuning surface properties in both organic photovoltaics and dye sensitized solar cells. This thesis explores the speciﬁc case of carboxylic acid modi-

fiers bonding to Zn1-xMgxO alloys. As discussed in Section 1.7, it is possible for these acidic modifiers to bond to the surface while simultaneously assisting in etching the surface. This is very important to carboxylic acid modification of Zn1-xMgxO.

The expected net result of carboxylic acid exposure to Zn1-xMgxO is that the carboxylate ion (deprotonated carboxylic acid) will coordinate to the Zn and Mg metal atoms to complete their coordination sphere.[81] This is in contrast to the TES attachment that bonds through Si to a surface hydroxyl oxygen molecule. Following coordination, the carboxylic acid will begin to etch the surface. In this work, two solvents were used to carry out the attachment process. Hexane was chosen for the study of benzoic acid modiﬁcation of Zn1-xMgxO discussed in Chp. 4. Hexane was chosen because it is a non-polar solvent and this should encourage the polar carboxyl attachment group to face the polar Zn1-xMgxO surface through electrostatic interactions.

Another advantage of the use of hexane is that it may reduce the etch rate of Zn1-xMgxO by benzoic acid by reducing the concentration of protons (H+) in solution. While most organic solvents are expected to perform much better than water in this regard, it is not known just how reduced the H+ concentration is compared to other solvents. Spalenka et al. compiled a list of proton concentrations of acetic acid in diﬀerent solvents for their study of stearic acid modiﬁcation of ZnO.[82] Unfortunately hexane was not included in their list, but the list

33 is helpful nonetheless. They found that THF produces the lowest H+ concentrations. We attempted to use THF as solvent in our studies but found that it adsorbs to the Zn1-xMgxO surface and produces IR modes in the carboxyl region of interest, signiﬁcantly complicating the background process. Ethanol was chosen as solvent for the dye sensitization studies in Chp. 5 since it is commonly used for dye sensitization of ZnO DSSCs.[70, 83]

2.4 Fabrication of Bulk Heterojunction Solar Cells

The performance of bulk heterojunction organic solar cells features prominently in Chp. 3 and their fabrication process will be discussed here. The full recipe for device fabrication is given in Chp. 3. This section will mainly describe the general deposition techniques and the purpose of each layer. The devices are deposited on a substrate of indium tin oxide on glass, which acts as a transparent conductor. The first step in the process is the deposition of ZnO according to the sol-gel process described in Section 2.1. The ZnO acts as the electron collecting contact, while simultaneously blocking holes from escaping the active layer. The second step is to modify the surface of the ZnO either through a cleaning process or by the molecular monolayer modification described in Sections 1.6 and 2.2 (Section 2.3 carboxylic acid modification was not used in our study of devices). This is not a necessary step for producing working devices, but Chp. 3 is based on the study of molecular modification of this contact. Next, the active layer must be deposited. In our devices, this was fabricated from a P3HT:PCBM blend. The deposition of this layer is itself a science. The key, as described in Section 1.4, is to create a very high surface area interface of P3HT and PCBM by forming a blend of these materials that is phase separated at a scale of ˜20 nm. The best process currently for accomplishing this is referred to as ’solvent annealing’.[84] The polymer:fullerene blend is first dissolved in a low volatility solvent such as dichlorobenzene, and then spin coated onto the substrate. Then the sample is enclosed in a small container, about the size of the sample, in order to trap solvent vapor. This dramatically increases the drying time of the film. The increased drying time allows the film enough time to phase separate to the optimal length scale. On top of the active layer, a hole

34 collecting layer is deposited. In our devices, this was a commercial variant on the conducting polymer poly(3,4-ethylenedioxythiophene):poly(styrenesulphonate) (PEDOT:PSS) that was spin coated from an aqueous solution and subsequently dried by thermal annealing. In our commercial material (Clevios HTL Solar, Heraeus), the composition has been tuned to allow the solution to wet the active layer effectively, possibly by the addition of surfactants. It also has substantially increased conductivity compared to standard PEDOT:PSS, again possibly a result of doping by surfactants. Finally, we deposited a silver top contact by thermal evaporation. To define the area of the device, a templated ITO substrate is used, while a mask is used to template the silver top contact. The two together define a square device. In testing our particular devices, we found the device must be masked while testing because the highly conductive hole transport layer makes lateral charge collection possible, defeating the careful definition of device area established in the fabrication process.

2.5 UV-Vis Absorption Spectroscopy

Ultraviolet-visible-near infrared (UV-Vis) absorption spectroscopy is used to characterize the optical response of a sample in these regions of the electromagnetic spectrum. The technique is based on the fundamental equation:

T (λ) + Abs(λ) + R(λ) = 1 (2.1) where T is transmittivity (or transmittance), Abs is absorptivity, R is reﬂectivity, and λ is wavelength. The absorbance (A, not to be confused with the absorptivity) of a sample is determined from its transmittivity as:

A = − log10(T ) (2.2) assuming that the sample’s reﬂectivity R is negligible or accounted for with an appropriate background. Absorbance, A, is the preferred quantity for describing optical absorption because it is directly proportional to the optical path length and the absorption coeﬃcient or molar absorptivity:

35 A = αl = εcl (2.3) where α is the absorption coefficient, l is the optical path length through the sample, c is the molar concentration of the absorber, and ε is the molar absorptivity. α and ε are independent of the sample dimensions but can depend on molecular orientation, anisotropy and other particular details of the material under study. The absorption coefficient, α is often used to describe the absorption of films, solids, pure liquids, or other pure substances. It can be used to calculate the thickness of a sample. The molar absorptivity, ε, is usually used to describe molecular absorption, where the absorbing molecule may be dissolved in a solution or dispersed over a surface. It can be used to calculate the concentration, c, of absorbing molecules if the path length is known. In this thesis, UV-Vis is mainly used to characterize a sample based on optically excited electronic transitions that result in absorption. For instance the absorbance spectrum of a semiconducting film can be used to calculate the material’s optical bandgap, Eg (Fig- ure 2.1(a)). Alternatively, the spectrum can be used to identify the energies of orbital- to-orbital transitions in molecular materials. For instance, the dye molecules used in dye sensitized solar cells undergo metal-to-ligand charge transfer (MLCT) transitions, and these account for the very strong visible absorption of these dyes. Shown in Figure 2.1(b) is the molar absorptivity of N3 dye (Figure 1.4) in ethanol and an absorbance spectrum of N3 on a ZnO surface. The spectrum of N3 on ZnO shows a chemical shift in the MLCT transitions that may be due to changing the chemical environment from ethanol to ZnO in air and/or to chemical bonding of the dye to Zn atoms. The absorption of N3 dye on Zn1-xMgxO is discussed extensively in Chp. 5.

2.6 Infrared Absorption Spectroscopy

Chemical structures (individual bonds, molecules, groups of atoms, or even crystalline solids) have characteristic vibrational frequencies that correspond to the infrared region of the electromagnetic spectrum. Infrared absorption spectroscopy measurements give information

36 Figure 2.1: UV-Vis absorbance spectra of materials used in this thesis. (a) Plot of absorbance squared used to find the bandgap (Eg) of Zn1-xMgxO films.[18] (b) Absorbance and molar absorbtivity of N3 dye (structure in Figure 1.4), with identification of orbital electronic transitions resulting in the absorbance spectra observed. MLCT = Metal to Ligand Charge Transfer. Note that density functional theory calculations suggest that the MLCT transitions actually occur from mixed Ruthenium-NCS orbitals and are thus not purely MCLT.[85] about the chemical structure of a sample through direct excitation of vibrational modes by absorption of infrared light. A vibrational mode of a chemical structure can be excited by infrared light only if the dipole moment of the structure changes during the vibration. Further discussion of this requirement can be found in Appendix A.1. The frequencies of vibrational modes are characteristic of the compounds in a particular material. The modes are normal modes of a set of coupled anharmonic oscillators made up of the atoms in the compound (masses) and their chemical bonds (springs). These normal modes can be approximately understood through the theory of coupled harmonic oscillators. While technically a compound is made up of a large number of coupled oscillators, often groups of atoms that vibrate at a characteristic frequency can be identified which oscillate relatively independently of the rest of the compound. The frequencies of these vibrations are called group frequencies.[86] Group frequencies are observable when a group of atoms vi- brates at a significantly different frequency from surrounding atomic groups, decoupling that

37 group from the rest of the compound.[86] Group frequencies generally fall within a certain limited range regardless of the compound the group appears in.[86] References compiling and discussing the vibrational frequencies of particular groups are available.[87] There are a number of ways to take advantage of group frequencies in infrared spectroscopy. In this thesis, group frequencies are mainly employed to:

1. Verify the presence of surface modifiers on a metal oxide surface through identification of particular vibrational modes in the spectrum that are expected based on the structure of the modifier

2. Analyze the conﬁguration of particular chemical groups through frequency shifts induced by the chemical environment of the bond

A Fourier transform infrared (FTIR) spectrometer has been employed for collection of the spectra in this thesis. The advantages and design of an FTIR spectrometer and discussion of how it operates is given in Appendix A.2. It is desirable that the contribution of ambient IR absorbers such as water and carbon dioxide be removed from the spectrum. In typical FTIR, this is accomplished with a purge of dry air or nitrogen that reduces water and carbon dioxide content and creates a stable atmosphere in which a background and sample can both be measured without significant changes to the ambient absorption. This process is time consuming and not entirely effective, especially when studying monolayers of material. The technique of polarization-modulation infrared reflectance-absorbance spectroscopy (PM-IRRAS) can overcome this disadvantage for certain carefully designed samples, and is used extensively in this thesis. In PM-IRRAS, the polarization of the incident infrared light is modulated (alternated) between parallel and perpendicular to the sample surface (Figure 2.2(a)) at a rate on the order of 100 kHz. If the sample is constructed on a metallic substrate and the sample is not too thick compared to the wavelength of the incident light, then the surface selectivity property of PM-IRRAS can be achieved. An example of such a sample is given in Figure 2.2(b), which

38 shows a thin ZnO film on gold that has attached to it a triethoxysilane surface modifier. The surface selectivity is accomplished through basic electrodynamics. The electric field of the parallel polarization (s-polarization, Figure 2.2(a)) will be cancelled by image charges formed in the conductor, while the field of the perpendicular polarization (p-polarization) will be enhanced by constructive interference. The absence of field strength parallel to the surface means that the sample will not absorb any s-polarized light. Therefore, any absorption by the s-polarized light will be due to background absorption, and the parallel polarization serves as an instantaneous measure of the background. Any infrared mode of the sample that has a component perpendicular to the surface will absorb p-polarized light. The p- polarization is therefore sensitive to both the background and the sample. The intensity of the s-polarized light (I(s)) can be subtracted from the p-polarized intensity (I(p)) to get the signal of the surface only (∆I = I(p)–I(s)). This is then divided by the sum I(p) + I(s) in order to get a spectrum that is independent of total light intensity. By measuring both the substrate and sample, an absorbance can then be calculated from the ratio of sample to substrate signal.[88]

2.7 Kelvin Probe Surface Potential Measurements

Surface potential measurements allow the determination of the difference in work function (Section 1.5) between two materials and estimates of the overall work function of a material. The basic working principle of Kelvin probe measurements is shown in Figure 2.3(a), while a simplified diagram of the Kelvin probe experimental setup is shown in Figure 2.3(b). A more in-depth discussion of how surface potential differences are determined using Kelvin probe can be found in Appendix A.3. The Kelvin probe experiment is set up by placing a conductive probe above a conducting or semiconducting sample (Figure 2.3(b)). In isolation,

probe and sample each have Fermi levels and work functions Ef,P and φP and Ef,S and φS, respectively (Figure 2.3(a,i)). The two materials are then connected electronically through the Kelvin probe system (Figure 2.3(a,ii)), and their Fermi levels equilibrate to form one constant Fermi level, Ef,Eq.. This happens via charge transfer and a potential diﬀerence (V )

39 Figure 2.2: (a) In polarization modulation infrared absorption spectroscopy (PM-IRRAS), the polarization of the infrared light is modulated between parallel and perpendicular to the sample surface. (b) Sample structure for PM-IRRAS measurements.

is established between the probe and sample. An electric field (E) due to this potential develops in the space between probe and sample. At equilibrium, V is referred to as the contact potential difference and is equal to the work function difference between the probe

and sample divided by e (V = (φS − φP )/e). The Kelvin probe system is setup so that an external voltage (Vappl) can be applied between probe and sample. This voltage can be tuned until V and E reach zero (Figure 2.3(a,iii)). At this point Vappl is equal and opposite to the equilibrium contact potential diﬀerence (φS − φP )/e. In this way, the relative work function φS − φP between probe and sample can be determined from Vappl. While this is the basic principle of Kelvin probe, there are some important details regarding determination of V = 0 and noise reduction techniques that are discussed in Appendix A.3. While the Kelvin probe can measure the relative work function between probe and sample, the work function of the probe itself is not known due to the sensitive electronics within the Kelvin probe system. Therefore, it is useful to ﬁrst measure a stable reference sample such

40 Figure 2.3: Basic working principle of Kelvin probe surface potential measurements. (a) Energy level diagrams showing the charge transfer processes occurring during Kelvin probe measurements. (i) In isolation the sample and probe have work functions φS and φP , respectively, with Fermi energies Ef,S and Ef,P . (ii) When the probe and sample are brought into electrical contact, their Fermi levels equilibrate through charge transfer to Ef,Eq., generating a potential drop V between probe and sample. (iii) Application of a bias voltage Vappl can be used to zero the electric field and achieve V = 0. Probe and sample will be in their equilibrium charge configurations. (b) Simplified schematic of the Kelvin probe experimental setup.

41 as gold and/or aluminum (with an intrinsic oxide layer) and then calculate the relative work function of the sample with respect to the reference.

2.8 Tapping Mode Atomic Force Microscopy

The basic goal of atomic force microscopy (AFM) is to gather information about the surface of a sample with nanoscale or even atomic resolution in both the lateral and vertical directions. While the original intent of the AFM was to gather morphological information (height vs. position) about the sample, a number of variants exist for measuring mechanical, chemical, electrical, and magnetic properties of the surface. The AFM measurements in this thesis focus on extracting morphological and mechanical properties of the surface of relatively fragile samples and therefore make use of tapping mode AFM (TM-AFM) and its associated phase contrast technique. The general technique of AFM is discussed in the introductory text by Haugstad.[89] In an AFM measurement of the morphology, a very sharp tip (usually with radius of curvature ∼10 nm) is rastered across the sample surface. In the simplest case, contact mode AFM Figure 2.4), the tip is put into direct contact with the surface. The tip interacts with the surface through local forces at the atomic scale. The tip is moved vertically away from or toward the sample by a feedback loop in order to maintain a constant force. The AFM is able to follow the morphology of the surface in this way. In tapping mode (Figure 2.4), also referred to as amplitude modulation AFM, the tip only makes periodic contact with the surface. This can prevent damage to fragile samples such as soft organic materials where a tip in contact mode could deform the organic materials. In tapping mode, the tip oscillates above the surface of the sample, making intermittent contact. It is driven just under its resonance frequency. Driving under resonance is understood to provide better images. Atomic forces interact with the tip over the course of an oscillation and these forces determine the amplitude of the oscillation by shifting the resonance frequency of the tip.[90] A feedback loop detects shifts in amplitude during scanning and adjusts the height of the tip in order to maintain a constant oscillation amplitude. A more detailed

42 Figure 2.4: Behavior of the tip in contact mode AFM (solid black line) and tapping mode AFM (dotted black line). In contact mode the tip is brought into contact with the surface and rastered across it while maintaining a constant force on the tip. In tapping mode the tip is oscillated at a higher average height above the sample, while making intermittent contact. The tip follows the topography by maintaining a constant oscillation amplitude.

explanation of how the AFM apparatus works in tapping mode can be found in Appendix A.4. The oscillating tip used in TM-AFM is sensitive to changes in the mechanical properties of the surface. These can be detected by a shift in the phase angle between the tip driving force and the tip amplitude. A tip driven at resonance in free space will always show a 90° phase lag of the amplitude behind the driving force (i.e. the force is at maximum when the tip amplitude is zero). However, in contact with a sample, this phase will shift. Moreover, the size of the phase shift depends on the mechanical properties of the sample. Speciﬁcally, the phase shift depends on the tip energy dissipated by the sample during a tip oscillation cycle.[90, 91] This idea is illustrated in Figure 2.5. A quantitative relationship between the phase and tip energy dissipation has been proposed[90, 91]:

ω A (ω) QE sin(θ) = sp + dis (2.4) ω0 A0 πkA0Asp(ω)

In this equation, θ is the phase angle, ω is the angular oscillation frequency, ω0 is the resonant

frequency, A0 is the resonant oscillation amplitude, Q is the tip quality factor, and k is the

tip spring constant. Asp(ω) is the amplitude at the scanning set point, maintained by the

feedback loop. Edis is the energy dissipated per oscillation cycle of the tip. Energy dissipation

43 calculated from measurement of tip force vs. distance curves agrees relatively well with this equation.[92]

Figure 2.5: Phase shift between AFM tip amplitude and driving force during TM-AFM. The figure shows a hypothetical flat sample made up of two different materials (blue and green) with significantly different hardness. The black sinusoid in the plots represents the driving force as a function of time. The blue and green sinusoids represent the amplitude of the tip as a function of time while scanning the blue and green materials, respectively. The phase shift is indicated in the plot.

Note that Equation 2.4 assumes that the feedback loop is able to maintain the amplitude

of the tip at Asp. However, feedback loops cannot respond instantaneously to changes in topography and the amplitude may deviate significantly from Asp when the topography changes quickly. It is apparent from Equation 2.4 that this will result in a phase shift due to a change in the first term. Stark et. al. have analyzed the proportional-integral feedback mechanism (A.4) and determined that the observed phase shift will be proportional to the first derivative of the height image, assuming that the contribution of the proportional component is small.[93] Figure 2.6 shows two examples of TM-AFM height and phase images. The scan of bare ZnO, produced through the sol-gel technique discussed in Section 2.1, shows how topography and phase are correlated. The edges of the ZnO grains in the height image, where the slope of the height image is steepest, are clearly seen in the phase image. The scan of octadecanethiol (ODT) on ZnO shows one of the advantages of phase images. ODT is composed of a long

44 alkyl chain with a thiol head group, and FTIR spectra of the sample shown in Figure 2.6 show that it is present on the surface of ZnO in multilayer coverage. Comparing the height image to that of bare ZnO it isn’t clear that there’s anything special happening on the surface, but the phase image shows clear contrast between the hard ZnO and the soft ODT multilayers. Similar phase images played a role in our analysis of bilayer ZnO/P3HT solar cells in which the ZnO layer was ﬁrst treated with ODT or octadecyltriethoxysilane.[55]

Figure 2.6: TM-AFM height and phase images showing the relationship between topography and phase and sample inhomogeneity and phase. The scale on the height is 20 nm while the scale on the phase is 1.5◦ on the bare ZnO image and 6◦ on the ODT on ZnO image.

2.9 X-Ray Photoelectron Spectroscopy

X-ray photoelectron spectroscopy (XPS) is a very surface sensitive measurement commonly employed in surface science. It is sensitive to approximately the ﬁrst 10 nm of

45 material.[94] In XPS a sample is loaded into ultrahigh vacuum and then bombarded with a monochromatic source of X-rays (Figure 2.7(a)). These X-rays penetrate deep into the sample (> 10 µm, Figure 2.7(b)) compared to the surface region of interest. Electrons in the sample absorb these X-rays and are ejected from their atoms. The electrons gain a kinetic energy (KE) equal to the diﬀerence in energy between the X-ray photon energy and their binding energy (BE) within the atom. While photoelectrons are generated as deep as the X-rays penetrate, the electrons of interest, those that avoid undergoing inelastic collisions, can only escape from a near-surface region of about 10 nm depth (Figure 2.7(c)). This is the process which gives XPS its very high surface sensitivity. Ejected photoelectrons are detected and sorted by BE using a KE-selective detector (Figure 2.7(d)). Since the KE energy of a photoelectron is determined by its BE under monochromatic X-ray irradiation, the BE of the photoelectron determines what depth it can escape from.[94]

The escape depth of a photoelectron is characterized by its inelastic mean free path (λIMFP ), which is the average distance traveled by a photoelectron before undergoing an inelastic collision. λIMFP ranges from a minimum of about 2-3 nm for photoelectrons of KE 20 eV to about 20-40 nm for electrons of KE 1000 eV (which is about the highest KE observed in an XPS experiment).[94] XPS spectra provide measurement of photoelectron signal intensity as a function of binding energy. Electrons in atoms have characteristic BEs depending on what orbital they reside in. XPS spectra can therefore be used to characterize what atoms are present in a sample and also in what proportion they appear in because of the characteristic orbital structure of the atom. Furthermore, the BE of a particular orbital shifts upon formation of a chemical bond by that atom. These shifts are on the order of a few eV and detectable in the line shape of the orbital peak. Therefore, the chemical state(s) of the atom in the sample can be determined by comparing the observed BEs to those of a reference. It should be made clear that the speciﬁc chemical bonds cannot be determined, just the chemical shifts in BE due

46 Figure 2.7: X-ray photoelectron spectroscopy (XPS) experimental setup. (a) A monochromatic source of x-rays is used to eject electrons from core orbitals. (b) X-rays can penetrate deep into a sample. However, photoelectrons can only escape from a thin layer (∼10 nm) near the surface without undergoing inelastic collisions (c). Electrons can escape from depths determined by their kinetic energy (KE). With monochromatic X-rays, this means that high binding energy (BE) electrons will have a thinner escape region than low BE electrons. Note that the electron escape and X-ray penetration regions are not drawn to scale. The photoelectrons are collected by a KE-selective detector which allows them to be sorted by binding energy.

47 to the bonding, which can be correlated with expectations for the bonding conﬁguration.

2.10 Contact Angle Goniometry

A drop of liquid on a solid surface will either spread out or bead up depending on the interaction between the drop and the surface. The extent to which the drop beads up can be quantiﬁed by the contact angle θc of the drop on the surface. The contact angle is the angle between the sample surface and the tangent to the water droplet edge at the water/surface boundary, shown in Figure 2.8(a). Contact angle measurements give useful information about the properties of the sample surface, including the surface energy. Changes in the contact angle indicate changes in the surface energy.[95]

Figure 2.8: (a) Contact angle measurement setup. A drop of liquid on a surface forms a contact angle θc according to the surface energy of the sample. (b) Examples of water contact angles on untreated, UV-Ozone cleaned, and monolayer modiﬁed ZnO surfaces.

Contact angles can be measured by placing a small droplet of liquid on the sample surface (Figure 2.8(a)). A diﬀuse light source is used to illuminate the drop and an image of the drop is captured using a camera. Spline ﬁtting software is then used to determine the contact

48 angle by fitting the shape of the droplet and determining the tangent angle at the edge of the drop. In this thesis, the DropSnake plugin for Image J was used to analyze contact angles.[96] Water contact angles of surface modified ZnO were measured in order to compare the effects of different modifications, including monolayer modification, on surface hydrophobicity. Figure 2.8(b) shows examples of the droplets observed. Generally, the contact angle increases after monolayer modification because of the hydrophobic nature of the organic functional groups. The functional group can have substantially different effects on the sample. The contact angle of PTES, which has a phenyl ring functional group, is substantially smaller than that of OTES which has an 18 carbon alkyl chain as functional group. The bare ZnO surface can display a large variety of CA angles depending on carbon contamination of the surface. A UV-Ozone cleaned surface will display a contact angle of around 18◦ due to hydroxylation of the surface and the lack of organic contaminants.[69] On the other hand a sample that has been allowed to sit in air or has been exposed to solvent is likely to show a higher contact angle. Consideration of this fact in the design of experimental controls is quite important.

2.11 Grazing Incidence X-Ray Diﬀraction

In crystalline atomic solids, atomic planes are spaced at distances that are on the order of the wavelength of X-rays (on the order of angstroms). Given the periodic arrangement of atoms, X-rays can reflect off atoms in one plane and constructively interfere with X-rays reflecting from other parallel planes. This occurs only if the path length difference between X-rays reflecting from the two planes is a multiple of their wavelength. This is known as the Bragg condition. This gives rise to certain angles of reflection showing much stronger signal through constructive interference. These angles of peak intensity are characteristic of the structure of the solid and can be used to gain a large amount of structural information.

In this thesis, structural information about Zn1-xMgxO thin ﬁlms produced by the sol-gel method of Section 2.1 is desired. The study of very thin metal oxide ﬁlms (in this case 25-50

49 nm) presents a challenge for obtaining structural information from X-ray diffraction. The penetration depth of X-rays is 10-100 µm into a sample.[97] For this reason, if conventional symmetric X-ray diffraction is used to study very thin films the dominant component of the signal will be from the substrate.[97] To isolate the signal from the sample, grazing incidence X-ray diffraction (GIXRD) uses an asymmetric measurement setup where the angle of incidence and reflection are not the same (Figure 2.9). The angle of incidence (δ in Figure 2.9) is chosen to be very small, close to the critical angle for total internal reflection, and is held fixed throughout the measurement.[97] This maximizes the path length inside the sample while minimizing the signal from the substrate. The detector is then rotated around the sample along the plane defined by the incident x-ray beam and sample normal, as shown in Figure 2.9. The angle formed by the detector line of sight and the surface is 2θ–δ, and the standard scattering angle 2θ can be obtained from this. Diffraction peaks in GIXRD measurements are found at comparable 2θ locations as in a symmetric XRD measurement.[97]

Figure 2.9: Experimental setup of grazing incidence X-ray diffraction. X-rays strike the surface at a grazing angle δ, which is chosen to be close to the critical angle for total internal reflection. δ is held fixed throughout the measurement. The detector makes an angle of 2θ–δ with the surface, and scans angles in the plane defined by the incident beam and sample normal.

50 2.12 Photoluminescence Spectroscopy

Photoluminescence (PL) spectroscopy probes the light emission spectrum of a sample after excitation with a light source. The excitation source excites electronic transitions from the electronic ground state to an excited state. This could involve excitation across the bandgap in a semiconductor or the MLCT transitions of N3 dye discussed in Section 2.5. After excitation, the electrons relax to the minimum energy of the excited state through fast non-radiative transitions. Without any other dominant mechanism of relaxation, the excited state will decay through emission back to the initial energy state through photon emission, usually on a much slower time scale than non-radiative relaxation. Due to the relaxation processes that occur prior to photon emission, the emitted photon will be at a lower wavelength than the exciting photon. The spectrum of this emission can be used to characterize the electronic structure of materials. The amount of PL obtained can also be an important topic of study. The PL signal can be quenched by non-radiative relaxation mechanisms if they dominate the decay process. Such quenching can occur when either the excited electron or hole is transferred to another material, as occurs in an organic or dye sensitized solar cell.[35, 98] Thus PL quenching can be a sensitive measurement of charge transfer. In ZnO-based dye senstized solar cells, dye can accumulate on the surface in Zn-dye complexes that do not inject charge efficiently into the ZnO substrate. These complexes fluoresce strongly at a characteristic wavelength identified by Horiuchi et al.[9] The PL signal from a dye-treated ZnO surface is therefore indicative of the amount of this complex on the surface. This technique is employed in Chp. 5 in the study of metal-dye complexes formed by the etching of Zn1-xMgxO films by N3 dye (Figure 1.4). The PL setup used in the experiments of this thesis is diagrammed in Figure 2.10(a). An argon ion laser is used to excite the sample. The 514.5 nm laser line was used in all measurements. The PL generated from the excitation is emitted in all directions. Some of this PL is collected by the collection optics and focused on the entrance slit of the spectrograph which

51 spatially separates the PL signal by wavelength. A long pass ﬁlter keeps scattered laser light from entering the spectrometer. The dispersed signal is detected by a CCD camera that records signal intensity as a function position on the CCD. The CCD pixels are correlated with wavelength allowing a spectrum of PL intensity vs. wavelength to be determined.

Figure 2.10: Illustration of photoluminescence spectroscopy experiment setup (a) and spectrum of N3 dye on ZnO (b).

Figure 2.10(b) shows the PL spectrum generated from a ZnO sample treated with N3 dye (Figure 1.4). The spectrum appears to be characteristic of metal-dye complexes that form on the ZnO surface, as discussed in Chp. 5.

52 CHAPTER 3 TUNING ZINC OXIDE/ORGANIC ENERGY LEVEL ALIGNMENT USING MIXED TRIETHOXYSILANE MONOLAYERS

Thomas M. Brenner 1, Gang Chen1, Erich P. Meinig 1, Darick J. Baker 1,*, Dana C. Olson2, Reuben T. Collins1, and Thomas E. Furtak 1

Originally published in J. Mater. Chem. C, 2013, 1, 5935. DOI: 10.1039/c3tc30881b. Available online. Reproduced by permission of The Royal Society of Chemistry.

Attributions: Thomas Brenner performed the PM-IRRAS, Kelvin probe, and contact angle measurements, interpreted the results, edited the paper, and wrote the bulk of it. Both Thomas and Gang Chen assisted in the production of ZnO thin films and their subsequent surface modification for all aspects of this work. Gang Chen built the photovoltaic devices and measured their J-V curves. Erich Meinig took atomic force microscopy measurements of the treated and untreated ZnO surfaces. Darick Baker first observed phenyltriethoxysilane in PM-IRRAS indicating its usefullness for this project. Abstract: Interfacial energy level alignment influences several critical organic optoelectronic device characteristics including charge transfer barriers, turn-on voltage, and open circuit voltage (V oc). Introduction of dipolar molecular monolayers on metal oxide surfaces has allowed improvements in device performance as well as fundamental studies of energy level alignment in these devices. We demonstrate that dipolar mixed monolayers can be covalently bonded to zinc oxide (ZnO) through the triethoxysilane chemical attachment scheme, and that these monolayers tune the work function of the ZnO surface over 0.6 eV. We then employ mixed monolayer-treated ZnO surfaces as the electron-selective contact in inverted

1Department of Physics, Colorado School of Mines, Golden, Colorado 80401, USA. 2National Renewable Energy Laboratory, Golden, Colorado 80401. *Now at Microfabrication Facility, University of Washington, Seattle, Washington 98195.

53 bulk heterojunction photovoltaic devices to determine the correlation between the V oc and the work function of the contact. We ﬁnd the relationship of -0.14 V/eV between the V oc and contact work function is consistent with current results and theories of contact inﬂuence on V oc.

3.1 Introduction

The issue of energy level alignment at interfaces is fundamental to the performance of organic electronic devices. Interfacial energy level alignment can determine, for instance, charge transfer barriers at contacts, the turn-on voltage of organic light emitting diodes

(OLEDs), and the open circuit voltage (Voc) of organic solar cells.[18, 49, 50, 52, 68, 99] Tuning the electronic structure of one or both materials in an interface, through several methods, has allowed energy level alignment at the interface to be studied and manipulated. This has led to both notable device improvements and results that further our understanding of how these devices and interfaces work.[18, 49, 50, 52, 57, 59, 68, 100–102] Metal oxides are widely applicable in organic electronics, and have proven to have highly tunable energy levels, making them amenable to energy level alignment tuning experiments. Moreover, the introduction of metal oxides into organic electronic devices has necessitated the study of metal oxide/organic interfaces and their energy level alignment behavior. Metal oxides can be degenerately doped to form transparent conducting oxide contacts such as indium tin oxide (ITO) or aluminum doped zinc oxide (AZO). They can be used as the acceptor in hybrid photovoltaic active layers, where their higher carrier mobilities could in principle boost efficiencies over organic-only devices. They can also be used as components of the recombination layer in tandem photovoltaic devices.[103–105] One of the most promising applications of metal oxides is as interfacial layers in bulk heterojunction (BHJ) photovoltaic devices. These interfacial layers are thin films (1-50nm) that separate a conductive contact from the organic active layer. Interfacial layers can have several functions in the cell, including controlling the compatibility of the contact with the organic, controlling energy barriers to charge transport between the active layer and contact, defining the polarity of a device,

54 and providing a contact with high charge selectivity.[51, 106] Both bulk and surface approaches have been used to tune metal oxide energy levels at an interface. One bulk tactic is to simply switch between different metal oxides.[59] Metal oxides have a variety of different electronic structures, ranging from n-type materials such as zinc oxide (ZnO) and titanium dioxide (TiO2) whose Fermi levels are closer to vacuum (lower work function–WF or φ) to materials with Fermi levels much further from the vacuum level (high φ) such as n-type molybdenum trioxide (MoO3) and tungsten trioxide (WO3) and p-type nickel oxide (NiO).[107] Another effective bulk approach is to tune energy levels by varying the composition of an alloy. For instance, ZnO can be alloyed with magnesium oxide (MgO) to tune its bandgap and WF.[18] Alternatively, the metal oxide surface can be altered. Surface treatments such as oxygen plasma or ultraviolet ozone cleaning can be used to alter the WF of ITO or ZnO.[61, 62, 65] However, there are reports that these treatments do not result in a stable surface.[64, 65] The WF can also be tuned by introducing monolayers of surface-bonded molecules with dipolar functional groups. Such dipolar monolayers were developed on metal and silicon surfaces and transferred to metal oxides.[76, 108–110] Many of these layers are stable over time and can produce large changes in φ (of the order 1eV).[64, 66] Dipolar functional groups introduce an electrostatic shift between the vacuum level inside the sample and the vacuum level just outside the surface, increasing or decreasing φ depending on the direction of the dipole.[111, 112] This change in φ (∆φ) can be determined from contact potential difference measurements as discussed in the experimental section. Figure 3.1(a) illustrates the effect of dipoles on interfacial energy level alignment at an idealized interface between a metal oxide and an organic. It is important to note that the dipoles introduced by a modifier originate both through the bond of the modifier to the surface as well as through the modifier’s functional group. Dipolar monolayer modification of metal oxides has been employed to demonstrate a number of important results in interfacial energy level alignment experiments. Introducing dipoles at the TiO2 surface through monolayer modification in hybrid TiO2/P3HT solar

55 Figure 3.1: (a) Energy level alignment at an ideal metal oxide/organic interface can be tuned by introducing a dipolar layer at the surface of the metal oxide before depositing the organic. The dipole introduces an electrostatic potential shift to charges crossing the interface, shifting the vacuum level between the metal oxide and organic (ELC ). The vacuum level shift will either decrease (top left) or increase (top right) the WF (φ) of the metal oxide surface compared to its unmodified state (top middle). The shift in work function (∆φ) can be measured as a contact potential difference. (b) Inverted bulk heterojunction (IBHJ) photovoltaic device (bottom middle) used to demonstrate the effects of tuning energy level alignment at the electron-selective contact/active layer interface. Dipolar monolayers (ML, bottom left and right molecules) can be inserted at the ZnO/P3HT:PCBM interface to shift energy level alignment.

56 cells has demonstrated that manipulating energy level alignment between the donor and acceptor directly affects Voc.[68] Many studies have focused on tuning the WF of a metal oxide contact. Tuning of the ITO contact in a conjugated polymer diode showed deviation from the Schottky-Mott limit.[113] Mixed results have been reported in attempts to change the Voc of organic bilayer solar cells or the turn-on voltage of OLEDs and organic diodes, with some studies showing a strong impact and others showing none at all.[99, 100, 114– 119] Interestingly, the study by Khodabakhsh et al. showed a strong dependence of solar cell short circuit current density (Jsc) on contact WF.[119] The study by Beaumont et al. suggests that whether Voc changes in bilayer small molecule devices depends on the choice of donor molecule, and the energy of its highest occupied molecular orbital (HOMO).[115] In BHJ solar cells, only a few studies of dipolar molecular modification of the contact to the active layer have been published. Ratcliff et al. demonstrated that Voc can be manipulated by tuning the WF of an ITO contact.[52] The dependence of Voc on contact WF was less than 0.3 V of Voc lost per 1 eV increase in WF (< 0.3 V/eV). The study also found that non-selective ITO contacts give consistently lower Voc than contacts that are expected to be selective such as NiO. Another studied employed alkylsilanes of differing terminal groups to tune the work function of ITO in contact with the BHJ active layer.[120] The largest work function sample showed the best performance in all of Voc, Jsc, and efficiency and this was attributed to both the contact work function change and surface energy changes from the hydrophobic terminal group. Monolayer modification of the metal oxide interlayer/metal contact interface in a BHJ device showed a strong dependence of device performance on the interface dipole.[121] ZnO is a promising metal oxide material for use in organic devices because of its compatibility with solution processing and the fact that its intrinsic n-type doping and band structure make it a good electron-selective layer.[3, 122] Several families of molecules representing a variety of attachment schemes have been used in molecular modifications of the surface of ZnO including: carboxylic acids, phosphonic acids, thiols, acetylacetonate, and

57 various silanes.[8, 54, 55, 67, 69, 99, 123–125] In this study, we have chosen to employ monolayers that attach via the triethoxysilane (TES) chemistry developed on ZnO by Allen et al.[69] The TES chemistry provides a strong covalent bond to the surface that shows no evidence of etching or multilayer formation as has sometimes been observed in common protic modifiers such as thiols and carboxylic and phosphonic acids.[8, 55] The ethoxy attachment group provides sufficient reduction in reactivity that the molecules do not readily cross-link or polymerize in the deposition solution as is found with other silanes.[69, 77] In this report, we verify that the TES modifiers phenyltriethoxysilane (PTES) and 4- chlorophenyl triethoxysilane (4CPTES) (Figure 3.1(b)) form conformal mixed monolayers on the surface of ZnO with relative concentrations on the surface comparable to their concentrations in the deposition solution. The functional groups of these modifiers have opposing dipoles that push the ZnO WF in opposing directions (Figure 3.1(b)). We demonstrate that the WF of the treated ZnO surface can be tuned by 0.6 eV by varying the composition of these mixed monolayers. This report is the first to demonstrate WF tuning of ZnO using a covalent attachment scheme (the TES scheme), and provides further verification of the viability of TES molecules as useful surface modifiers. To explore the effect that tuning the ZnO WF has on devices, we employ our modified ZnO as an electron-selective contact in inverted bulk heterojunction (IBHJ) solar cells with a poly(3-hexylthiophene) (P3HT) : phenyl-C61-butyric acid methyl ester (PCBM) active layer (Figure 3.1(b)). The open circuit voltage (Voc) of the device increases as the ZnO WF decreases, with a slope of -0.14 V/eV. This result is discussed in the context of the current understanding of how the WF of the contact affects the Voc of organic solar cells.

3.2 Experimental

ZnO Preparation, Surface Functionalization, and IBHJ PV Device Construction: 40 nm polycrystalline ZnO layers were deposited on glass microscope slides (Fischer Scientiﬁc), indium tin oxide (ITO) patterned glass (Thin Film Devices) and gold-on-glass substrates (Platypus Technologies) by the following sol-gel method: A solution of 0.75 M zinc acetate

58 dihydrate and 0.75 M ethanolamine in 2-methoxyethanol was stirred at 800 rpm on a hot plate at 60◦C for 30 min. After the solution cooled, it was spin coated onto a substrate at 2000 rpm for 60 s. The sample was then annealed at 300◦C for 10 min. in ambient air. This method is based on that of Ohyama et al., with modiﬁcation as discussed by Olson et al.[53, 62, 73] To prepare ZnO surfaces treated with molecules attaching through the triethoxysilane

(TES) chemistry, the ZnO samples were ﬁrst rinsed with toluene and acetone, dried with N2 and UV-ozone cleaned for 15 min. to induce surface hydroxylation. The samples were then placed in a solution of TES molecules at 45◦C for 90 min. The solution consisted of 0.072 M n-butylamine and a total of 0.18 M combined of phenyltriethoxysilane (PTES) and/or 4-chlorophenyltriethoxysilane (4CPTES), in toluene. To vary the ratio of PTES to 4CPTES molecules on the surface, the molar ratio of PTES and 4CPTES was varied in solution. The molar ratios used were 1:0 PTES:4CPTES (all PTES), 2:1 (2/3 PTES, 1/3 4CPTES), 1:2, and

0:1. After soaking, the samples were rinsed with toluene and acetone, blown dry with N2 and placed in a 110◦C oven for 60 min. After baking, samples were rinsed with toluene and acetone and dried with N2. This procedure follows that of Allen et al. for attaching TES molecules to ZnO.[69] Control samples for the TES treatment were prepared by following the procedures above, except that the molecular modiﬁers were excluded from the deposition solution. These controls were used as references for PM-IRRAS, Kelvin Probe, and CA measurements and were used to build control IBHJ devices. Devices for these experiments were fabricated on patterned ITO substrates that were sonicated in acetone and isopropyl alcohol for 10 min. each, followed by a 30 min. UV Ozone treatment. A 40 nm thick ZnO layer was deposited on the ITO by the sol-gel method described above. The ZnO surface was then treated before active layer deposition. For the PTES/4CPTES treated and control devices, the ZnO layer went through the surface functionalization steps as described above and then was rinsed with toluene and acetone,

◦ dried with N2, followed by 5 min. annealing in air at 100 C. A standard device was fabricated

59 following an already published procedure for preparing this type of device.[122] In this device, the ZnO ﬁlm was rinsed with DI water, acetone and isopropyl alcohol, dried with N2 and then annealed at 200◦C for 5 min. After the treatment of the ZnO surface, the active layer for all devices was spin-coated from a 50 g L-1 solution of P3HT (Rieke Metals) and PCBM (Nano-C) in a 1:1 ratio by weight in anhydrous dichlorobenzene at 800 RPM with 800 RPM s-1 acceleration for 60 s in a glove box. The wet ﬁlms were then allowed to dry slowly for

90 min. in covered petri dishes in N2 atmosphere.[84] This yielded active layers of thickness ∼245nm. CleviosTM HTL Solar PEDOT:PSS (Heraeus) was spin-coated on top of the active layer at 2000 RPM for 60 s in ambient air, followed by 5 min. annealing at 100◦C in ambient air and subsequent 10 min. annealing at 120◦C in a glove box. Silver was thermally evaporated onto the PEDOT:PSS at a rate of 0.5 A˚ s-1 until 30 nm was reached and at 1 A˚

-1 ◦ s for another 70 nm. Completed devices were annealed in N2 atmosphere at 130 C for 10 min. ZnO Surface and IBHJ PV Device Characterization: Polarization Modulation - InfraRed Reflectance Absorbance Spectroscopy (PM-IRRAS) measurements were performed to verify the presence of functional molecules on the ZnO surface, and determine the ratio of PTES to 4CPTES molecules on the surface. PM-IRRAS provides an extremely surface-sensitive IR absorbance measurement by cancelling environmental noise through polarization modulation of the incident light.[126] Measurements were performed on samples on gold-on-glass substrates using a Thermo Scientific Nicolet 6700 FT-IR spectrometer with a Nexus PEM Module. Spectra were taken between 650-4000 cm-1. For all samples, the ZnO surface was measured before and after surface modification. The post-modification measurements were divided by the pre-modification measurements to generate a spectrum proportional to absorbance.[88] Contact angle (CA) measurements were performed using sessile drops of deionized water. The CA images produced were analyzed using the DropSnake algorithm.[96] Atomic Force Microscopy images were taken in tapping mode using a Bruker Nanoscope III instrument

60 with high-aspect ratio 125µm cantilever tips (Nanoworld NCSI). Both topographic and phase information were collected. Roughness analysis of the topography images was performed using Nanoscope 5.30r1 software. Differences in WF (∆φ) between surface treatments were determined by contact potential difference (CPD) measurements taken in air, in the dark, for samples on gold substrates using a KP Technology SKP5050 Kelvin Probe. The CPD of each treatment was measured with respect to a gold standard. However, the ∆φ are reported in reference to the bare ZnO control surface described above, assigned a value of ∆φ = 0 meV. Devices were characterized under an AM1.5 simulated solar spectrum with 100 mW cm-2 illumination intensity and in the dark in an inert environment. An aperture was not used during measurement, so current densities and efficiencies are slightly higher than would be expected. Current density-voltage (J-V) curves were measured using a Keithley 236 source meter. Series resistance was calculated from the slope of the J-V curve at open circuit voltage. Light soaking was performed under the same light source by placing the devices under illumination for 20 min. During the light-soaking and subsequent measurements, devices were not temperature-controlled.

3.3 Results and Discussion

ZnO thin films prepared from a sol-gel precursor were treated with PTES, 4CPTES, and two solutions containing mixtures of PTES and 4CPTES molecules as discussed in the experimental section above. Figure 3.2(a) shows the spectrum, for each treatment, of polarization modulation infrared reflection absorption spectroscopy (PM-IRRAS) measurements in the phenyl ring vibration region. The legend indicates the ratio of PTES (P) to 4CPTES (4C) based on their molar concentrations in each deposition solution. Also shown is the spectrum of the control (explained in experimental section), which was prepared identically to the treated samples except no surface modifiers were included in its soaking solution. Based on the structure of each molecule, the major peaks shown were identified. The peak at 1090 cm-1 was identified as a vibration of the phenyl ring combined with a carbon-chlorine (C-Cl)

61 stretch characteristic of 4CPTES only. The peak near 1130 cm-1 was identified as a ring vibration combined with a carbon-silicon (C-Si) stretch.[87] This peak, though present in both PTES and 4CPTES treated samples, falls in intensity with increasing 4CPTES content. The peak is smaller in 4CPTES because the highly electronegative chlorine atom takes electron density away from the atoms involved in the 1130 cm-1 mode, thereby reducing the IR activity of the mode. The presence of these peaks in the spectra of treated samples, but not in that of the control, demonstrates that these peaks can be used as an indicator of the presence of each molecule on the surface. The increase in the 1090 cm-1 peak seen with increasing 4CPTES concentration in the deposition solution and the drop in the intensity of the peak at 1130 cm-1 indicates that the proportion of each molecule on the surface is changing as its proportion changes in solution. This qualitative observation can be made quantitative by using the integrated intensities of these peaks to calculate the fraction of each modifier on the surface. The results of these calculations are shown in Figure 3.2(b), where the calculated fraction of 4CPTES is plotted against the fraction of 4CPTES in the deposition solution. The details of this calculation are given in the supplementary information (Appendix B.1). The fraction of 4CPTES on the surface increases monotonically with its fraction in solution, with a slope close to one. This demonstrates that the fraction on the surface is comparable to the fraction in solution, within experimental error. Water contact angle measurements provided further verification that the molecules were present on the surface. All treatments showed consistently larger contact angles than for the deposition control by at least 10◦. These values are also at least 20◦ higher than the contact angle for a UV-Ozone cleaned ZnO surface, which has been measured in previous work.[69] In addition, the contact angle fell by 8◦ as the proportion of 4CPTES increased, indicating that changing the proportion of the modifiers on the surface directly affected the surface energy. The average contact angle measurement for each sample and its uncertainty are given in the supporting information (Appendix B.2). Atomic force microscopy (AFM) scans showed no difference in roughness between the treated and untreated surfaces. Differ-

62 Figure 3.2: (a) PM-IRRAS measurements of ZnO films treated with PTES and 4CPTES in different proportions. The legend indicates the molecular ratios in the deposition solutions based on molar concentration, with P = PTES and 4C = 4CPTES. Two peaks identified as vibrations of the phenyl ring in each molecule are shown. (b) The fraction of 4CPTES on the surface of ZnO based on integrated intensities from the PM-IRRAS measurements of (a), plotted against the fraction of 4CPTES in the deposition solution.

63 ences in roughness would be expected if multilayer or island-like agglomerates were present on the treated surface. Phase images also showed a homogenous surface in contrast to observations we have made with molecular treatments that leave agglomerated material or multilayers on the sample.[55] AFM phase images and a table of roughnesses can be found in the supporting information (Appendix B.3). Our measurements of infrared spectroscopy, water contact angle, and atomic force microscopy in combination with previous work on this chemistry provide strong evidence that there is no multilayer or aggregate formation and instead a monolayer has formed on the surface.[55, 69] It should be noted that this layer is expected to be lower in coverage than a close-packed alkyl chain monolayer.[69] The possib- lity of these molecules cross-linking on the surface may limit the minimum distance between them, reducing the overall coverage and packing.[69] Alternatively, the low coverage also suggests possible monolayer defects such as pinholes or uncovered regions.[127] There may also be segregation of the two molecular species involved, forming domains of PTES and 4CPTES.[128] These defects may have a significant impact on interface electronic properties and on the device as a whole, as discussed below.[127] Kelvin probe contact potential difference (CPD) measurements of treated ZnO relative to the bare control were used to determine the differences in work function (∆φ) between the various treatments. These measurements showed that varying the composition of the deposition solution allowed φ to be systematically tuned over 0.6 eV (Figure 3.3). As the proportion of 4CPTES increased, the WF of the surface increased as well. This is in direct agreement with our expectations for φ based on the dipole direction of each molecule. In principle, any φ between that of PTES and 4CPTES can be achieved by choosing the correct proportion of the two. It should be noted that the Kelvin probe measures a macroscopic average CPD, and therefore does not detect microscopic variations due to, for instance, segregation of the two molecular species.[128] We also note that the relationship between CPD and the composition of a mixed monolayer may not be linear. One interesting unexplored question suggested by our study is whether mixed monolayers might show the non-linear depolarization effects seen

64 when changing the coverage of a single species monolayer.[112, 127]

Figure 3.3: Contact potential diﬀerence (relative work function) ∆φ of treated ZnO referenced to the bare control surface as a function of PTES and 4CPTES mole fraction in solution. ∆φ increases monotonically with increasing 4CPTES proportion, as expected from the diﬀerent dipole moments of the two molecules. The results presented are representative, and the error bars indicate the reproducibility of a measurement in a particular experiment.

In order to understand how tuning φ through monolayer modification would affect the properties of a ZnO/organic interface we prepared inverted bulk heterojunction photovoltaic devices with the structure shown in Figure 3.1(b). In these devices, ZnO acts as an electron- selective contact (ESC), extracting electrons from PCBM and blocking transport of holes in either P3HT or PCBM. We modified the surface of the ZnO with mixed PTES and 4CPTES molecular layers before spin coating the P3HT:PCBM blend. The ratio of PTES to 4CPTES was varied using the same ratios as used in the PM-IRRAS and Kelvin probe measurements. It is important to note that we also constructed two different untreated devices: a standard device and a control device. As discussed above, the control device contains a ZnO surface treated in the same way as the TES modified surfaces, but without including the TES modifiers in the deposition solution. This was the control surface used throughout the PM-

65 IRRAS, Kelvin probe, contact angle, and AFM measurements. The ZnO in the standard device is treated according to a previously published procedure for making IBHJ devices with bare ZnO interlayers.[122] Both the control and standard device processing procedures are given in the experimental section. The ZnO surface in the standard device is optimized for device efficiency, while the control surface is designed to compare surfaces treated identically with the exception of the presence or absence of TES molecular modifiers. The standard and control devices give insight into what effect the surface preparation procedures for TES deposition have on the device performance. Figure 3.4 shows representative current density - voltage (J-V) curves under illumination for these devices after 20 min. light soaking, while Table 3.1 gives their photovoltaic parameters. Dark J-V curves can be found in the supplemental information (Appendix B.4). Light soaking was necessary because the devices initially showed a double diode or s-kink (see supplemental information, Appendix B.5). Double diodes are commonly observed in the literature, but are not yet well understood.[129– 133] Light soaking is an established method of minimizing the influence of the double diode on performance.[132, 133]

Table 3.1: Characteristics of IBHJ devices with monolayer modified ZnO ESCs after 20 min. light soaking. FF = fill factor, η = power conversion efficiency, RS = series resistance, RSH = shunt resistance. A mask was not used during device testing so the high conductivity of the PEDOT:PSS layer may have enlarged the collection area of the devices, leading to apparent current densities that are too large.

2 2 2 Treatment Voc [mV] Jsc [mA/cm ] FF η [%] RS [Ωcm ] RSH [Ωcm ] Standard 591 11.13 0.48 3.14 10.0 225 Control 565 11.20 0.44 2.80 12.9 190 1:0 P:4C 532 10.83 0.45 2.62 15.5 231 2:1 P:4C 504 10.68 0.41 2.23 17.2 168 1:2 P:4C 492 9.95 0.44 2.15 17.8 235 0:1 P:4C 451 10.70 0.39 1.90 18.1 153

The short circuit current (Jsc) showed no trend (over multiple experiments) with respect

to modiﬁer proportion. However, Jsc was always less in treated devices compared to untreated devices. From the standpoint of surface energies, the reduced current is counter to

66 Figure 3.4: Representative light J-V measurements of IBHJ photovoltaic devices containing mixed monolayer modified ZnO. Jsc of treated devices shows no trend while Voc falls by 100 mV as the concentration of 4CPTES increases. what might be expected since the surface energy decreases with treatment, and this should improve the blend morphology near the interface, which has been shown to increase Jsc.[125] Some explanations for the lower current are that the modifier molecules may increase surface recombination or traps or introduce a small barrier to charge extraction at the interface. A charge extraction barrier would be consistent with the increased series resistance of treated devices, but its effects may go beyond just resistive losses.

The open circuit voltage (Voc) of the treated devices increases systematically with PTES concentration, increasing on average by 95±12 mV in transitioning from 4CPTES to PTES.

While the Voc of all the devices was improved by light soaking, this Voc trend was not affected by light soaking. This indicates that the trend is due to the dipole introduced by the surface modifiers on ZnO. The surface dipole can be understood to affect Voc by changing the energy level alignment between the contact and the active layer. As the work function of the contact decreases, electrons can be extracted at higher and higher energies, increasing

Voc. This explanation is supported by computational modeling of bulk heterojunction solar

67 cells.[43, 134] The variability in ﬁll factor (FF, Table 3.1) between treatments suggested that

there may be recombination processes occuring that could also impact Voc. However, the FF

variation can be explained by its correlation with shunt resistance (RSH ) according to the

equivalent circuit model of the solar cell. RSH does not correlate with Voc in our devices and for this reason cannot contribute to the trend in Voc, but may contribute to Voc overall. Several other details in the device results are important. First, all of the TES treated devices, as well as the control, have a lower Voc than the standard device preparation. The reason for this is likely the UV-Ozone (UVO) clean used during TES attachment and in preparing the control. This is done to hydroxylate the surface which promotes TES attachment.[69] UVO cleaning of ZnO thin ﬁlms, however, has been shown to induce a dipole unfavorable to electron collection, and other changes to electronic structure.[16, 62] We also note that the workfunction of the untreated control surface falls between that of the 4CPTES and

PTES treated surfaces (Figure 3.3) while the Voc of the treated surfaces is always less than for the untreated. This discrepancy can be explained by the fact that the relative surface dipole between the treated and untreated samples could shift significantly between Kelvin probe measurement and within the active layer, since contact with air (for Kelvin probe measurement) and deposition of the active layer could affect the treated and untreated samples quite differently. It is also possible that the monolayer introduces barriers or traps, leading to increased surface recombination and lower Voc. This explanation is in agreement with the lower Jsc and fill factors observed in treated devices.

Perhaps of more significance is the observation that while the change in Voc with respect to ∆φ for the mixed monolayer treatments is linear, the slope is -0.14 V/eV (Figure 3.5), with the negative sign indicating that increases in ∆φ decrease Voc. This result is of interest because the influence of contact WF on Voc in BHJ devices is not well established, especially for situations in which monolayer modification is involved. In organic bilayer devices, the majority of studies suggest that the Voc is essentially uninfluenced by the WF of the contact, and device modeling agrees with these results.[43, 115, 118, 119] For BHJ devices, a range

68 of relationships between the contact work function and Voc have been proposed and demonstrated. Some studies have seen a relationship of less than 0.3 V/eV which is similar to

the result in Figure 3.5.[49, 52] The study by Mihailetchi et al. suggests Voc should depend strongly on the contact work function when injection barriers are high, while it should show a small dependence once the contact work function approaches the HOMO or LUMO of the organic.[38] This relationship is supported by the experimental work of Tress et al., which shows that the slope of Voc vs. the HOMO of a hole extraction layer (HHEL) decreases as HHEL approaches the HOMO of the donor.[43] Our results are consistent with the model of Mihailetchi et al. and the experimental work of Tress et al. since we observe a shallow relationship of -0.14 V/eV while tuning the WF of ZnO, which is near the LUMO of PCBM.[51] The study by Tress et al. also provides a model showing that BHJ devices should show a dependence of Voc on the contact work function. However, the dependence is 1 V/eV, which does not agree with our results or experimental results in general.[43] One possible reason for this is that any charge transfer or reorganization that might occur at the interface is not taken into account in the model. The study by Mihailetchi et al. suggests a low slope is due to charge transfer from the contact into a state in the organic which pins the Fermi level of the contact.[38] This explanation is supported by metal oxide and metal/organic interface studies showing such Fermi level pinning at the HOMO or LUMO.[57, 59] Of course, other factors may influence this slope as well. For instance, the density of states (DOS) of the donor and acceptor have a significant impact on Voc in the active layer, and the DOS may also affect Voc through contact energy level alignment as well.[40] The shallow slope of Voc vs. WF may also be explained in our particular case by the quality of our monolayer, something that is not often considered in analyses of monolayer modified organic electronic devices. Scanning tunneling microscopy (STM) and ballistic electron emission microscopy (BEEM) measurements of Au/Dipolar Monolayer/GaAs interfaces suggests pinholes in the monolayer and the surrounding monolayer domains each give rise to separate Schottky barriers.[135] Further experiments suggested that current is funneled through the pinholes and that the

69 pinholes determine the behavior of the interface.[127, 136] Large pinholes can completely cancel the effect of a dipolar monolayer, while smaller pinholes are affected by the fringing fields of the monolayer. Because our molecular layers are expected to be sub-monolayer in coverage, device characteristics may be affected by pinholes that are not detectable in our

macroscopic CPD measurements, and these pinholes could lead to a low Voc vs. WF slope. This may be an issue in other interface modiﬁcation studies as well. The heterogeneity introduced by a mixed monolayer might impact interface electronic properties as well. Charge transfer may be favored through one species over another and in this case interface properties may not change much with composition. Interestingly, this does not appear to be the case in our results since the mixed monolayers ﬁt the trend between the homogeneous endpoints.

More investigation is necessary to determine the relationship between contact WF and Voc, especially in the case of mixed and single species monolayer modiﬁcation at the interface.

Figure 3.5: Plot of Voc for devices with the indicated surface treatment of the ZnO electron- selective contact against the relative work function ∆φ of the ZnO surface with the same treatment. The slope is -0.14 V/eV.

70 3.4 Conclusions

We have successfully attached mixed monolayers of covalently bonded triethoxysilane dipolar surface modifiers to ZnO, as verified by several surface analysis methods. Using these modifiers, we can shift the work function of ZnO over 0.6 eV demonstrating the viability of triethoxysilane as a solution-deposited surface attachment scheme for tuning the work function of metal oxides used in organic electronics applications. We used our monolayer- modified ZnO as an electron-selective layer in IBHJ photovoltaic devices. We demonstrated that the Voc changes systematically as a function of contact work function, in agreement with current models and experimental results for contacts to BHJ organic solar cells.

3.5 Acknowledgements

The authors would like to thank the National Renewable Energy Laboratory (NREL) for access to the Kelvin probe and the organic photovoltaics lab. T. M. B. and G. C. would like to thank Edwin Widjonarko for Kelvin probe training and Sarah Cowan for device fabrication training. T. M. B. would like to thank Jao van de Lagemaat for helpful discussion. This material is based upon work supported by the National Science Foundation through Grant No. DMR-0907409 and through the Renewable Energy Materials Research Science and Engineering Center under Grant No. DMR-0820518. D. C. O.’s funding is provided by the DOE EERE SETP program.

71 CHAPTER 4

ETCH-RESISTANT ZN1-XMGXO ALLOYS: AN ALTERNATIVE TO ZNO FOR CARBOXYLIC ACID SURFACE MODIFICATION

Thomas M. Brenner 1, Thomas A. Flores2,*, Paul F. Ndione3, Erich P. Meinig 1, Gang Chen1, Dana C. Olson3, Thomas E. Furtak 1, and Reuben T. Collins1

Attributions: Thomas Brenner performed the bulk of the infrared and UV-Vis spectroscopy measurements, analysis, and interpretation, and wrote the paper. Thomas Flores took the etch rate measurements and performed infrared and UV-Vis spectroscopy on bare

Zn1-xMgxO samples and samples exposed to BA for 1 hr. Paul Ndione provided the grazing incidence X-ray diffraction measurements and interpretation. Erich Meinig took AFM measurements of the bare Zn1-xMgxO films. Gang Chen helped develop our Zn1-xMgxO deposition process and provided training and guidance for the undergraduate authors. Abstract: Zinc oxide (ZnO) is a high bandgap transparent metal oxide used extensively in organic electronic devices and dye-sensitized solar cells (DSSCs) as an electron accepting and transporting material. Carboxylic and other organic acid attachment schemes are often employed to sensitize or functionalize the surface of ZnO in these applications. Unfortunately, one weakness of ZnO is its high susceptibility to etching by even weak acids. This has a substantial negative impact on ZnO DSSCs which use carboxylic acid groups to attach dyes to ZnO, and also influences attachment techniques for functionalizing ZnO in organic electronic

1Department of Physics, Colorado School of Mines, Golden CO 80401 2Department of Physics, Lehigh University, Bethlehem PA 18015 3National Renewable Energy Laboratory, Golden CO 80401 *Now in the Department of Applied Physics, Stanford University, Stanford CA 94305

72 devices. By substituting Mg for Zn atoms to form Zn1-xMgxO (ZnMgO) alloys, a material with similar electronic properties but higher etch resistance is achieved. Here, we show that the etch rate of Zn0.8Mg0.2O, when exposed to the prototypical modifier benzoic acid (BA), is an order of magnitude lower than that of ZnO. Infrared spectroscopic characterization of BA-modified ZnMgO indicates that BA binds to the surface in two different conformations, the relative proportion of which changes with Mg content. The IR spectra also provide evidence of the presence of physisorbed zinc benzoate complexes that are likely a product of etching, as their presence increases with exposure time. Possibilities for the mechanism of improved etch resistance are discussed via the literature on metal oxide dissolution.

4.1 Introduction

ZnO is a high bandgap transparent metal oxide used extensively in organic electronic devices and dye-sensitized solar cells as an electron accepting and transporting material. The two characteristics that make ZnO a good electron acceptor are that it is intrinsically n-type and its conduction band is well matched with the LUMO of many organic semiconductors and sensitizing dyes.[3, 51, 137] The deep valence band of ZnO makes it an effective hole blocking material as well.[51] In organic electronic devices, organic/inorganic interfaces and their properties are critical to performance. However, ZnO/organic interfaces are not necessarily optimal, and the properties of the interface can be improved through surface modification of the ZnO. Monolayers of surface-bonded molecules have proven effective in modifying the surface properties of ZnO and thus the properties of ZnO/Organic interfaces.[54, 69, 79, 99, 138, 139] However, some promising modification schemes have been shown to produce non- ideal monolayers through acid etching of the ZnO surface under long deposition times.[8, 140] The problem of etching has also been identified in dye-sensitized solar cells (DSSCs) employing dyes with acidic attachment groups on ZnO. These dyes dissolve the ZnO surface and form Zn-Dye aggregates that absorb light efficiently but transfer charge poorly.[9, 37, 70, 83, 141] These aggregates could be referred to as ’photon parasites’. The optimal soaking time in ZnO-based DSSCs is therefore a balance of dye infiltration into the nanostructured

73 ZnO matrix and etching of already sensitized regions.[70, 83] ZnO-based DSSC devices could be improved if etching could be reduced. The process by which highly adsorbing acids such as the above organic surface modifiers etch metal oxides has been studied previously and informs this work. In general, dissolution of metal oxides by acids (commonly referred to as etching in much of the literature) proceeds by three steps: protonation of the surface oxo groups (metal-oxygen-metal or Me- O-Me groups), rupture of the Me-O bonds, and phase transfer of the Me from surface to solution.[71, 72] In the case of semiconducting metal oxides, the acid anion can play a significant role in dissolution through complexation of surface metal atoms.[71, 72] Under very strong adsorption, complexation of the surface Me atoms by the adsorbing species is assumed to isolate these atoms from the rest of the solid, and the dissolution chemistry is dominated by these complexes.[71] In this case the rate-limiting step is understood to be the transfer of this surface complex to solution.[72] Because metal oxide surface modification by organic acids is based on adsorption of the acid anion on metal oxide surfaces, the dissolution behavior is likely to approach this case. Garcia-Rodenas et al. have been able to formulate a theory explaining the rates of dissolution of simple (single metal) semiconducting metal oxides by a highly adsorbing organic acid.[72] This is discussed below in the context of our results. The issue of etching by organic acid surface modifiers has been approached in several ways. Highly dynamic deposition processes (such as spin coating) have been employed to create modifier layers from organic acids in organic electronic devices. Such an approach likely limits exposure to the acid and can minimize etching.[99, 121, 138, 142] The results of these processes may be more difficult to reproduce because highly dynamic processes are more difficult to control than equilibrium processes. The results of our work suggest more investigation is needed to confirm the effectiveness of this approach in preventing etching.

Another approach to reducing etching is to consider the acid dissociation constant (pKa) of organic acids in diﬀerent solvents. It has been found that acetic acid shows a pKa of 24 in

74 tetrahydrofuran compared to 10.3 in ethanol, indicating a drop in proton concentration of several orders of magnitude.[82] Thus, deposition from tetrahyrdofuran will result in reduced etching compared to ethanol. This is an eﬀective approach if the desired attachment is

soluble in high pKa solvents. In DSSCs, the problem of Zn-dye aggregate formation has been approached through the design of less acidic or non-aggregating dyes or by covering the ZnO in a thin layer of etch-resistant material.[37, 143–145] To overcome the issue of etching of ZnO by acidic molecular modifiers, we explore the use of Zn1-xMgxO (ZnMgO) alloys. Such alloys have been previously used as an electron acceptor in hybrid photovoltaic cells.[18] This study demonstrated that the work function of the alloy decreases as Mg content is increased, and this can be employed to raise the open circuit voltage of the device and improve efficiency. The study suggests that ZnMgO alloys are suited for use as an electron transport or accepting layer throughout organic electronics. The main drawback of the material is that its conductivity decreases exponentially as Mg content increases.[17, 18, 48] The improved etch resistance of ZnMgO alloys to acids was first observed by Taratula et al. while studying acid pre-treatments of ZnO and ZnMgO films in preparation for surface modification. They found that exposure to acid at pH < 4 could be accomplished with films of ZnMgO (5-10 %Mg) whereas ZnO films were etched.[8] This etch resistance may reduce the formation of aggregates or etch products and make possible longer depositions of organic acid modifiers that would normally etch ZnO, leading to more reproducible surfaces and potentially to better monolayers and improved organic/inorganic interface control. Here we quantify the etch resistance of solution-derived ZnMgO thin films in the presence of the archetypal carboxylic acid modifier benzoic acid (BA), and demonstrate that the etch rate for alloys containing 20% Mg is an order of magnitude smaller than that of ZnO. Our infrared spectroscopy measurements of the alloys soaked in BA provide evidence that benzoic acid is attached to the surface in two distinct bonding conformations the relative proportion of which changes with Mg content. Furthermore, we provide evidence of the formation of zinc benzoate complexes during the etching process that

75 subsequently physisorb to the surface. We hypothesize a mechanism for the improved etch resistance based on our ﬁndings and on existing literature about dissolution of semiconducting metal oxides.

4.2 Experimental

Zn1-xMgxO (ZnMgO) thin films were deposited on microscope slide glass (Fisher Scien- tific) and gold-on-glass (50 nm Au, 2.5 nm Ti adhesion layer, aluminosilicate glass, Platypus Technologies) substrates through the following sol-gel process: A total molar concentration of 0.75M zinc acetate dihydrate and magnesium acetate tetrahydrate was dissolved in a solution of 2-methoxyethanol and 0.75M ethanolamine. Zinc acetate and magnesium acetate were combined in the desired molar ratio (1-x and x, resp., where 0 < x < 0.3) to achieve the 0.75M total concentration. The solution was stirred at 800 rpm on a 60◦C hot plate for 30 min. After the solution cooled, it was spin coated onto a substrate at 2000 rpm for 60 s. The sample was then annealed at 300◦C for 10 min. in ambient air (except where stated otherwise) in order to decompose the zinc and magnesium acetate and allow them to react with oxygen. The films were cooled with flowing nitrogen and then placed on a metal cooling plate before being used in experiments. All films were created from sol-gel solutions that were not more than 1 day old in order to avoid the formation of precipitates. This method is based on the procedure by Olson et al.[18] A Multimode atomic force microscope (AFM) with a Nanoscope III controller from Bruker was used to characterize the morphology of the untreated ZnMgO alloy surfaces. Root mean square (RMS) roughness was calculated from the collected morphology data. The ZnMgO thin films were treated with BA either by spin coating for 60 s at 2000 rpm from a 2 mM solution in hexane, or by immersion in a 2 mM solution in hexane. Hexane was chosen as solvent to encourage the polar carboxylic acid attachment group to orient toward the film surface as opposed to into the non-polar solvent. After treatment, the films were rinsed with pure hexane and dried with flowing nitrogen for profilometer and infrared spectroscopy measurements. After measurement, the samples were re-immersed in the BA

76 solution if more treatment was desired. UV-vis absorption spectroscopy was performed on the untreated and BA treated ZnMgO films on glass slides using a Cary V spectrophotometer. The spectra were taken between 800 and 200 nm and were used to determine the bandgap of the material. The crystalline structure of the films was determined by θ − 2θ X-ray diffraction (XRD)

measurements using a Rigaku D/Max-2500 diffractometer with CuKα radiation (λ = 1.5418A).˚ In order to characterize these very thin films (20-50nm), X-ray diffraction (XRD) data were collected at a grazing angle (of 0.3◦) between the X-ray beam and the film surface to maximize the signal ratio from the ZnMgO films compared to that from the glass. To measure the thickness of the ZnMgO films for etch rate calculations, a corner of the film was removed from the substrate using 1 M hydrochloric acid. Profilometer measurements of the sample edge were then taken using a Tencor P-10 Surface Profiler. Measurements of the same edge were taken during repeated exposure of the films to the BA solution. Polarization Modulated - InfraRed Reflectance Absorbance Spectroscopy (PM-IRRAS) measurements were performed on the untreated and BA treated ZnMgO films. PM-IRRAS provides an extremely surface-sensitive IR absorbance measurement by minimizing environmental noise and emphasizing surface-localized species through polarization modulation of the incident light.[126, 146] Measurements were performed on samples on gold-on-glass substrates using a Thermo Scientific Nicolet 6700 FT-IR spectrometer with a Nexus PEM Module. Spectra were taken between 650-4000 cm-1. Each alloy was measured before and after BA treatment. The spectra of the BA treated surfaces were divided by spectra from their respective untreated surfaces to generate an absorbance spectrum.[88] Peak fits in the carboxyl asymmetric stretch spectral range were performed using Omnic 8.0 FTIR analysis software. Fits were performed by first fitting the spectra from the least-etched samples - those containing 20 or 30% Mg. These fits yielded two peaks with peak positions near 1550 and 1570 cm-1. The variability in the peak position and full-width-half-maximum (FWHM), as defined by the observed range of these quantities in the fits to the least-etched samples,

77 was then employed to define bounds on these quantities to be used in fits to the remaining spectra (0 and 10% Mg samples). In these more highly etched samples, new peaks were introduced if those derived from the 20 and 30% Mg spectra were insufficient to produce a good fit.

4.3 Results and Discussion

Figure 4.1 shows bandgaps determined from UV-VIS measurements of our films as a function of the Mg content of the sol-gel solution from which they were deposited. The bandgaps were calculated according to Ref. 18, which is based on the theory by Urbach.[147] The bandgap increases with increasing Mg content, indicating an alloy has been formed. This is in agreement with the results of Refs. 17, 18, 74. We conclude from the bandgap variation that we were able to produce ZnMgO alloys of varying Mg content. The error bars Figure 4.1 reflect variation from one experiment to the next. While all films were prepared in the same way, the bandgap of 30% Mg content films showed significant variability, and this will be discussed further below. The bare ZnMgO samples were also characterized with grazing incidence X-ray diffraction (XRD) in order to verify their expected crystal structure, determine their crystallinity, and determine any trends related to Mg content. While ZnMgO prepared by the sol-gel method discussed above has been characterized before, the variability in bandgap we observed in our 30% Mg films suggested that characterization of our particular films could be insightful.[18] Figure 4.2 shows the background-subtracted XRD patterns of ZnMgO thin films at different Mg concentrations. The films are identified as having a typical wurtzite ZnO structure with no other phases, indicating these films are polycrystalline and single phase in the entire composition range under study. All diffraction peaks are indexed according to the diffraction data of hexagonal wurtzite structure ZnO (JCPDS #361451). XRD peaks are observed at 2θ = 31.9◦, 34.3◦, 36.1◦. It is noticed that the (002) XRD peak intensity becomes weaker and the (002)/(101) peak intensity ratio becomes smaller as the Mg content increases, indicating an alteration in the crystalline quality. Mg content doesn’t have much influence on the

78 Figure 4.1: Bandgaps of ZnMgO films as a function of Mg content (black points), as determined from UV-VIS measurements. Significant variation was observed in the bandgap of samples made with 30% Mg which is reflected in the large error bar. Our measured bandgaps are in good agreement with those measured by Olson et al. in Ref. 18 (red points). peak positions except for 30% Mg-doped films where the (100) peak position moves toward lower angles (31.7◦), demonstrating an increase of the a-axis, which could be related to the presence of compensated defects and/or residual strain.[148, 149] Two 30% Mg content films were considered in our XRD study, one with high bandgap (film ’a’ in Figure 4.2), consistent with the measurements of Olson et. al. (Figure 4.1), and one with low bandgap (film ’b’).[18] Film ’b’ exhibits low peak intensity and higher noise compared to the other films presented here, including film ’a’, indicating a lower degree of crystallinity. Tapping mode AFM measurements showed that low bandgap films of 30% Mg showed appreciably higher RMS surface roughness than the lower Mg content alloys. The variability observed for 30% Mg content samples in bandgap, crystallinity, and surface roughness indicates that it is difficult to control the film formation process with our preparation method when Mg content is 30%. One possible explanation is that phase segregation occurs, forming separate domains of wurtzite ZnMgO and rock salt MgO of unknown compostion, as has previously been observed.[17, 18, 74, 75, 148, 150] Ohtomo et al. found

79 Figure 4.2: Background-subtracted grazing incidence X-ray diffraction spectra of ZnMgO with 0, 10, and 30% Mg content. Two films prepared with 30% Mg content (films ’a’ and ’b’) are shown. Both these films were characterized because they showed substantially different bandgaps (indicated in the figure) even though they were prepared identically. Film ’b’ is of substantially lower crystallinity, as evidenced by its lower signal counts and high noise.

80 that incorporation of Mg into the rock salt MgO structure and not the wurtzite structure reduces the bandgap of the film, since the wurtzite structure has lower Mg content and therefore a lower bandgap.[74] Our XRD results do not directly support this explanation because no rock salt peaks are observed in the spectra. However, the very low crystallinity of the low bandgap sample suggests there may be amorphous material present of unknown composition that may be affecting the optical properties of the wurtzite ZnMgO present. The variability in film properties observed for 30% Mg alloys is likely indicative of a high sensitivity to the processing conditions. This is supported by previous studies that conclude ZnMgO alloys are metastable and the maximum achievable Mg content depends on the processing conditions.[74, 75, 148] Another issue with Mg containing films was the presence of unreacted magnesium acetate, identified through IR spectroscopy. This did not impact surface treatment appreciably, as discussed below. When our ZnMgO films were soaked in BA, they displayed lower etch rates with increasing Mg content, indicating enhanced etch resistance in the alloy vs. ZnO. The etch rate was calculated by comparing film layer thicknesses measured by profilometry after various soaking times to the original layer thickness. The etch rates for ZnMgO alloys of varying Mg content are given in Table 4.1. The table shows that the etch rate of alloys containing 10% Mg is one quarter that of ZnO and alloys with 20% Mg have an etch rate one order of magnitude smaller than ZnO. The etch rates of 20 and 30 % Mg content films were indistinguishable.

Table 4.1: Etch rate of ZnMgO films soaked in 2 mM BA in hexane for different Mg concentrations as observed through profilometer measurements taken during exposure of film to BA solution.

Material Etch Rate (nm/min.) ZnO 0.29±0.06 Zn0.9Mg0.1O 0.07±0.04 Zn0.8Mg0.2O 0.03±0.05 Zn0.7Mg0.3O 0.03±0.04

81 In order to determine the presence and binding configuration of BA on the surface of ZnMgO films, PM-IRRAS measurements were performed. These measurements provide evidence of BA binding to the ZnMgO film and evidence of the accumulation of etch products in low Mg content samples. Spectra were taken of ZnMgO films treated with BA in the following ways (as discussed in the experimental section): by spin coating, soaking for 30 min. or soaking for 1 hr. in a 2 mM solution in hexane (Figure 4.3). It should be mentioned that a unique experiment was performed for each treatment: a separate sample was used for each of the exposure times. The spectral region shown in Figure 4.3 contains the symmetric

− − - (νsym(CO2 )) and asymmetric (νasym(CO2 )) stretch modes of the CO2 carboxylate bonding group. Also found in this region are several vibrations of the phenyl ring (RM = Ring Mode). The vibrational modes associated with each feature were identiﬁed through reference to an infrared study of zirconium dioxide (ZrO2) treated with BA as well as an infrared study of zinc benzoate complexes.[81, 151] Fits were performed for all spectra as described in the experimental section, allowing the IR features to be resolved into peaks associated with

− specific vibrational modes. Figure 4.4 shows the peaks present in the νasym(CO2 ) feature most critical to our analysis of BA surface bonding, from samples soaked in BA for 1hr. The fits in Figure 4.4 are representative of all the spectra, and demonstrate all of the asymmetric stretch features observed in this study except the Figure 4.4 mode observed on the high energy side of the ring mode at 1600 cm-1, near 1639 cm-1 (discussed below, see Figure 4.3). The peaks identified in the fits, their mode assignment and association with any surface species are given in Table 4.2.

− Table 4.2 indicates there are five distinct νasym(CO2 ) modes indicating at least five bonding configurations of BA. Each of these is discussed below. We find that, for a given Mg content, the peaks near 1552 and 1571 cm-1 (red peaks in Figure 4.4) are present with relatively constant integrated intensity across exposure time. This leads us to the conclusion that these peaks are associated with BA bonded to the ZnMgO surface. It is also apparent from Figure 4.4 that the relative intensity of these peaks varies systematically with Mg

82 Figure 4.3: PM-IRRAS infrared absorption spectra of BA-treated ZnMgO films showing the carboxyl stretch region. The spectra are sorted by exposure method/time. Features are identified as either a vibration of the phenyl ring (ring mode = RM) or a symmetric - − or asymmetric stretch of the CO2 carboxyl group (νsym/asym(CO2 )). There are multiple − νsym/asym(CO2 ) modes of differing origin in the spectra and they are only generically labeled here.

83 − Figure 4.4: Peak ﬁts to the dominant νasym(CO2 ) feature of the IR spectra of ZnMgO ﬁlms soaked in BA for 1 hr. Red peaks = BA bonded the ZnMgO surface. Green peak = etch product (normal zinc benzoate). Blue peak = etch product (basic zinc benzoate). See main text and Ref. 151.

Table 4.2: Peaks identified in fits of the IR spectra of BA treated ZnMgO. The peaks − are sorted by the mode to which they correspond. νsym(CO2 ) = symmetric stretch of − carboxylate group, νasym(CO2 ) = asymmetric stretch of carboxylate group, RM = ring mode - vibration of phenyl ring. The surface species associated with a particular peak is noted where identification was possible. Surface species marked with * were identified through Ref. 151.

Mode Approx. Peak Position (cm-1) Surface Species 1410 ν (CO−) sym 2 1420 1450 RM 1495 1600, 1620 1532 Etch Product* 1552 Surface Bond − νasym(CO2 ) 1569 Etch Product* (0% Mg) 1571 Surface Bond (>0% Mg) 1639 Etch Product*

84 content. This will be discussed further below. The study of zinc benzoates by Clegg et al.

− provides an interpretation for the νasym(CO2 ) peaks near 1530 (green peak in Figure 4.4) and 1639 cm-1 (Figure 4.3).[151] These peaks are only detectable in low Mg content ﬁlms exposed to BA for a long time (0 & 10% Mg in Figure 4.4). Given the low etch resistance of ZnO, we suggest that these peaks are associated with a product of etching that has exceeded its solubility limit in the solvent and is physisorbed to the surface. In fact, these peaks correspond quite well to those observed in the ’normal’ zinc benzoate complex.[151] This complex is a polymeric structure in which BA has two distinct bidendate bridge-bonded

-1 − conﬁgurations.[151, 152] The mode at 1639 cm , which is a very high frequency νasym(CO2 ) mode, is due to an unusual bridging geometry.[151, 152] The formation of this complex is similar to the complexes formed through etching of the ZnO surface by dyes in DSSCs.[9, 70, 141]

− The above discussion focuses on the ﬁne structure in the dominant νasym(CO2 ) fea-

− − ture. The splitting of the νsym(CO2 ) and νasym(CO2 ) modes (∆) is also of interest and often indicative of the manner in which a carboxylate species is bonded.[153] Unfortunately,

− ﬁne structure of the νsym(CO2 ) feature was not discernible in our ﬁts because of the close

− − proximity of the νsym(CO2 ) modes. This prevents a unique pairing of the νasym(CO2 )

− − modes with their corresponding νsym(CO2 ) modes. Depending on how the νasym(CO2 )

-1 − modes of the surface-bonded species at 1552 and 1571 cm are paired with the νsym(CO2 ) modes, ∆ is between 130-170 cm-1 suggesting a bidendate bonding conformation of BA to the surface.[81, 153] While the assignment of bidentate binding has a high probability of being correct due to the much larger ∆ expected for monodentate binding, determining whether the bonding falls under bridging or chelating is difficult without corroborating evidence from other measurements. Further conclusions about the bonding configurations of the modes at 1552 and 1571 cm-1 can be made based on the variation of peak intensities with Mg content of the films. Carboxylic acids adsorb to both ZnO and MgO at room temperature, creating a coordination

85 to the metal.[99, 121, 138, 142, 154, 155] We therefore suggest that BA will coordinate to both the Zn and Mg atoms on the surface of ZnMgO. Given that the vibrational modes of BA bonded to Zn versus Mg are diﬀerent, this will result in a systematic variation of the relative

− integrated intensity of the νasym(CO2 ) modes of BA as Mg content changes. This can be observed qualitatively in Figure 4.3 and Figure 4.4; the peak at 1570 cm-1 grows progressively taller with Mg content, relative to the peak at 1550 cm-1. Figure 4.5(a) demonstrates this quantitatively by showing that the ratio of the integrated intensity of the 1570 cm-1 peak to that of the 1550 cm-1 peak increases monotonically with Mg content, regardless of exposure time. This indicates a bimodal shift in bonding conﬁguration from one distinct mode to another as Mg content increases. Based on these trends, our interpretation is that the peak near 1550 cm-1 corresponds to BA bonded to Zn while the peak near 1570 cm-1 corresponds to BA bonded to Mg.

− Figure 4.5: (a) Plot of the ratio of the integrated intensity of the νasym(CO2 ) mode at 1570 cm-1 to that at 1550 cm-1 as a function of Mg content for each exposure time. The trend is toward a higher peak ratio with greater Mg content, indicating a shift in integrated intensity from the lower to higher wavenumber peak as Mg content increases. (b) The same peak ratio, plotted against bandgap.

86 There are two qualifying factors associated with this tentative assignment that we will now address. The first is that we must assume BA does not bind to both Zn and Mg simultaneously (in a bridging conformation) because this would produce another feature in the IR spectrum that is not observed. This could be the case if the bonding configuration is different between the two metals (i.e. bidentate chelating in Mg and bridging for Zn) or is chelating for both. The second issue is the apparent presence of the peak near 1570 cm-1, assigned to BA bonding to Mg, in ZnO samples that contain no Mg (blue peak in Figure 4.4). This peak is strongest in ZnO soaked for 1 hr. and decreases in intensity as exposure time decreases (Figure 4.6), unlike the peak in higher Mg content samples which has constant intensity with exposure time. This trend correlates well with a decrease in the normal zinc benzoate etch product (Figure 4.6). This indicates that the peak near 1570 cm-1 in the ZnO spectra is likely also due to a physisorbed etch product with a mode energy that overlaps that

− of BA bonded to Mg. We suggest that the identity of this peak is actually the νasym(CO2 ) mode of another zinc benzoate complex, termed the ’basic’ complex, whose peak location has previously been reported to be 1562 cm-1.[151] The basic complex can be synthesized from the normal complex by re-crystallization, and it is possible that this conversion process occurs at high concentrations of the normal complex or at all concentrations but at a slower rate than the formation of the normal complex.[151] The basic complex has a tetrahedral Zn4O core coordinated to six benzoate ions. The benzoate ions have only one bidentate bridging

− − conformation and therefore only one νasym(CO2 ) mode.[151] Since the νasym(CO2 ) mode of the basic complex is expected to appear near 1562 cm-1, it can be considered co-located with the mode we have assigned to BA bonded to Mg in higher Mg content samples, but actually has a diﬀerent origin and is just coincidentally located at nearly the same position. We also note from Figure 4.6 that spin coating appears to create a substantial amount of etch product, counter to the reasoning behind its employment as a surface modiﬁcation technique. The conditions in spin coating are very dynamic and it is possible that very high concentrations of BA are achieved in the solution as the solvent dries, leading to strong

87 etching.

− -1 Figure 4.6: Variation of the νasym(CO2 ) modes at 1532 and 1569 cm in ZnO (0% Mg) samples as a function of exposure time to BA. As exposure time increases, the integrated intensity of these peaks also increases, suggesting that they correspond to etch products that − accumulate on the surface over time. These peaks have been assigned as the νasym(CO2 ) modes of normal (1532 cm-1) and basic (1569 cm-1) zinc benzoate complexes.[151] Note the presence of the normal zinc benzoate complex in spin coated samples.

We have also pursued an alternative analysis of these results exploring correlations between the IR features and bandgap instead of Mg content. We took this alternative approach because of the variation we observed in the bandgap and other propreties of samples containing 30% Mg. Figure 4.5(b) indicates that the relative proportion of the two surface bonded species can be equally well explained by the variation in the bandgap of the ﬁlms, which may be indicative of the amount of Mg actually incorporated into the alloy.[74] The most interesting result from our analysis using bandgap as the independent variable is shown in Figure 4.7, which gives the integrated intensity of each of the 1550 and 1570 cm-1 peaks and their sum as a function of bandgap. The integrated intensity of both peaks and their sum increases with bandgap suggesting that surface coverage increases with bandgap. This is a result that deserves further investigation because the increased surface coverage should create more dramatic changes in surface properties. These results can be explained by as-

88 suming that the processes of etching and surface attachment compete to achieve a steady state. Etch resistance increases with Mg incorporation into the alloy, for which we consider bandgap a proxy. As etch resistance increases, the steady state will shift away from etching by benzoic acid toward higher surface bonding of benzoate. These conclusions are consistent with those made by Baumeler et al. while studying sensitization and etching of ZnO by dyes of varying acidity and with the fact that the etch rate is understood to be limited by the rate of transfer of surface complexes to solution.[72, 141]

-1 − Figure 4.7: Plot of integrated intensity of 1550 and 1571 cm νasym(CO2 ) modes and their sum as a function of bandgap.

We also observed the presence of unreacted magnesium acetate in the IR measurements of our films (Figure 4.8(a)). Unreacted acetate was identified through PM-IRRAS by its symmetric and asymmetric stretch modes in the bare films.[156] This fact has not been previously reported and raises concerns about whether surface modification of the films might be affected by the presence of this acetate. To test if unreacted acetate influenced our surface modification study, a series of Zn0.8Mg0.2O films were created using different temperatures for the annealing (decomposition) step described in the experimental section. The amount of unreacted acetate decreased as decomposition temperature increased, as measured by the

89 integrated intensity of the carboxyl region in PM-IRRAS (Figure 4.8(b), horizontal axis). The ﬁlms were then treated with BA for 30 min. as discussed in the experimental section,

− and PM-IRRAS was measured again. The integrated intensity of the νasym(CO2 ) feature of BA bonded to the ZnMgO surface (identified above) was used to characterize the amount of BA attached (Figure 4.8(b), vertical axis). We conclude from this experiment that the unreacted acetate does not play a significant role in surface functionalization of ZnMgO films by BA because the amount of BA attached to the surface does not change significantly with the amount of acetate. At decomposition temperatures higher than 380◦C, the integrated

− intensity of the BA νasym(CO2 ) feature falls dramatically. This is attributed to changes in crystallite orientation expected under such high temperature anneals.[73]

Figure 4.8: Relationship between amount of unreacted acetate in ZnMgO films and amount of attached BA (a) PM-IRRAS absorbance curve of an untreated Zn0.8Mg0.2O film dissociated at 315◦C. The background for this spectra is bare gold, so the spectra shows the absorbance − − of the film itself. The carboxyl region is shown, and the νsym(CO2 ) and νasym(CO2 ) modes of unreacted magnesium acetate are the dominant features. (b) The amount of unreacted acetate in the film decreases with increasing decomposition temperature. The plot shows that the amount of BA attached to the surface (vertical axis) stays roughly constant while the amount of unreacted acetate in the film (horizontal axis) decreases by a factor of four between dissociation temps. of 250 − 380◦C. At high annealing temperatures, it has been reported that the crystallite orientation changes dramatically and this could be the cause of the drop in BA attached to the surface.[73]

90 Finally, we provide some insight into the possible origins of etch resistance in ZnMgO based on current literature of acid dissolution of metal oxides. One important consideration in the etching process is the point of zero charge (pzc). The point of zero charge is the point at which surface hydroxide and surface proton concentrations are balanced.[71] This occurs

at diﬀerent values of pH (referred to as pHpzc or pH0) for diﬀerent materials, and therefore determines the balance of hydroxide and protons at a particular pH. The concentration of surface protons is one factor that will determine the etch rate; the more adsorbed protons, the higher the etch rate. Therefore, all else being equal, it is expected that a material with

higher pHpzc will etch faster than a material with a lower pHpzc at a given pH. ZnO has pHpzc of around 9, meaning that under acidic conditions the surface will have a large imbalance of protons vs hydroxide.[70] Part of the etch resistance of ZnMgO may arise from a lowering of

the pHpzc upon addition of Mg. However, the pzc is not the deﬁning variable in the etch rate.

A number of very stable oxides (such as Al2O3 and NiO) also have pHpzc of 9 or above.[71] Alternatively, the reactivity of the metal oxide lattice itself can be considered. A theory explaining the dissolution rates of simple (single metal) metal oxide semiconductors in highly adsorbing media can be found in a study of ZnO, NiO, CoO, and Fe2O3 dissolved in oxalic acid.[72] This study suggests the dissolution rate is limited by the rate of transfer of metal- oxalate complexes from the surface to solution, and that the rate of transfer is determined by how easy it is to break the bond between the metal complex and surface oxygen (the lability of the bond). The lability of the bond is a property of the metal that can be

estimated from the water dissociation rate from the metal ion in water (k−w).[72, 157] k−w is essentially the rate at which water, acting as a ligand for the ion, escapes its coordination to the ion.[157, 158] The slower this rate, the less labile the metal ion is. Rodenas et al.

have argued that the correlation of k−w with the rate of dissolution of simple metal oxides by oxalic acid indicates that the etch rate is determined by the lability of the ion.[72] The

2+ 2+ 2+ lability of Mg should be substantially less than Zn since the k−w of Mg is two orders of magnitude smaller than that of Zn2+.[158] It may be the case that a stronger bond of Mg to

91 the metal oxide surface improves the overall stability of the compound. However, as is often the case in mixed metal oxides, we might expect a leaching effect in which the more reactive Zn2+ ions are selectively dissolved from the surface, leaving behind a MgO scaffold. This is an important aspect to consider, as the dissolution rate can be limited by the leach rate and leaching changes the overall structure and composition of the surface.[71, 72] If indeed the lability of the metal ion, and not a leaching effect, is determining the dissolution rate, we can

2+ suggest another alloy that may make further improvements to etch resistance. Ni has k−w

+ an order of magnitude lower than Mg2 .[158] Thus the compound Zn1-xNixO with wurtzite crystal structure may have improved etch resistance compared to ZnO and Zn1-xMgxO. This compound has been synthesized and its high-bandgap semiconductor properties established, but its dissolution properties have not.[75, 159]

4.4 Conclusions

ZnMgO oxide alloys show significant improvements in etch resistance against benzoic acid compared to ZnO, with alloys containing 20% Mg showing an order of magnitude drop in etch rate. The reduction in etch rate could allow substantially longer, more controllable deposition of carboxylic or phosphonic acid surface modifiers in organic electronic devices or dye sensitized solar cells. We have also completed the first infrared characterization of benzoic acid-treated ZnO and ZnMgO alloy films, and identified the carboxyl stretching vibrations of the benzoate ion bonded to the surface. Benzoic acid bonds in two different configurations to ZnMgO, and the relative proportion of these configurations varies systematically with Mg content. We have also provided infrared evidence that suggests that zinc benzoate complexes are produced during the etching process and are physisorbed to the surface.

4.5 Acknowledgements

This material is based upon work supported by the National Science Foundation through Grant No. DMR-0907409 and through the Renewable Energy Materials Research Science and Engineering Center under Grant No. DMR-0820518. DCO’s funding is provided by the

92 U.S. Department of Energy under Contract No. DE-AC36-08-GO28308 with the National Renewable Energy Laboratory through the DOE SETP program.

93 CHAPTER 5

EXPLORING THE MECHANISM OF ZN1-XMGXO ETCH RESISTANCE THROUGH DYE SENSITIZATION

Thomas M. Brenner 1, K. Xerxes Steirer2, Erich P. Meinig 1, Dana C. Olson2, Thomas E. Furtak 1, and Reuben T. Collins1

Attributions: The majority of sample preparations and measurements were performed by Thomas Brenner. Erich Meinig determined the molar absorptivity of N3 dye and made an initial exploration of dye sensitization of ZnMgO, demonstrating the feasibility of the project. The XPS measurements of ZnMgO modified by benzoic acid were performed by Xerxes Steirer. Abstract: Surface modification of metal oxides with molecular monolayers is an effective strategy for tuning surface and interface properties in excitonic devices employing metal oxide acceptor and transport layers. Monolayers are used to sensitize metal oxides in dye sensitized solar cells (DSSCs) and to optimize interfacial properties in both organic and DSSCs. The most commonly used attachment chemistries are acid/base reactions employing organic acids. The use of acid/base chemistries has presented a problem for one of the most commonly used and promising metal oxides in excitonic devices, ZnO. Unfortunately, ZnO is easily etched by even weak organic acids, and this can lead to non-ideal monolayers and the accumulation of surface Zn complexes formed in the etching process. This is especially troublesome in ZnO-based dye sensitized solar cells where etching leads to the formation of Zn-dye complexes that reduce the incident photo to current conversion efficiency. In previous work (Chp. 4), Zn1-xMgxO (ZnMgO) alloys have been considered as an alternative to ZnO displaying higher resistance to acid etching. ZnMgO alloys displayed a 10-fold reduction in

1Department of Physics, Colorado School of Mines, Golden CO 80401 2National Renewable Energy Laboratory, Golden CO 80401

94 etch rate when exposed to the organic acid benzoic acid, a prototypical surface modifier, as compared to ZnO. These results suggest that ZnMgO is a promising alternative that may alleviate some of the problems with ZnO discussed above. However, questions about the mechanism of etch resistance were posed in the previous work that are important to consider for future applications of the material. In that work, a long-time, or steady-state, etch rate was measured during which significant film loss occurred. It was suggested that the etch rate could be limited by compositional and morphological features of the surface established only over long time scales of etching. However, for applications of this material as a substrate for dye sensitization or interface tuning, the initial, or transient etch rate is really the quantity of interest. Thus, the questions of what the mechanism of etch resistance is in ZnMgO and what, if any, reduction in the initial etch rate this affords are key for the applications of this material. In order to explore these questions, we investigated the sensitization of ZnMgO by N3 dye (cis-bis(isothiocyanato)bis(2,2’-bipyridyl-4,4’-dicarboxylato)-ruthenium(II)), which attaches through carboxylic acid groups. The relatively slow etch rate of N3 dye that we observed, combined with its strong optical absorption made it possible to study the initial etch rate of the ZnMgO film. From observations of the rate of accumulation of metal- dye aggregates on the surface, we conclude that the initial etch rate of ZnMgO increases with Mg content, in contrast to the steady-state etch rates observed for ZnMgO exposed to benzoic acid. Photoluminescence and PM-IRRAS measurements of the N3 sensitized ZnMgO surfaces show that the etch products produced are primarily Zn-N3 complexes regardless of Mg content. XPS measurements of benzoic acid modified ZnMgO support the conclusion that in general Zn-carboxylate complexes are the major products of etching. In order to try to engineer the surface composition and morphological changes hypothesized to limit the steady-state etch rate of ZnMgO, the films were briefly exposed to a dilute mineral acid before sensitization. The accumulation of Zn-dye complexes on pre-etched ZnMgO surfaces showed a minimum at 5-10% Mg content instead of the monotonic increase observed for as-deposited films. This suggests that a pre-etch may be an effective technique for limiting Zn-dye complex

95 accumulation. We conclude that our results are consistent with an etch resistance mechanism that is manifested only after a transient period during which the surface is altered in order to establish a Mg-rich ZnMgO scaffold layer. Based on the identification of surface complexes, we propose that Zn is selectively leached from the surface leaving behind an Mg-rich scaffold structure at the surface which is resistant to further etching and also protects subsurface ZnMgO by limiting the rate of diffusion of etch species toward and away from the subsurface region.

5.1 Introduction

The field of dye sensitized solar cells, after more than a decade of stagnation, is making progress again, reaching a world record efficiency of 12.3% in 2012.[2] Breakthroughs in dye design, electrolyte design, and strategies for light trapping and suppressing recombination have made improvements possible.[2] As discussed in Section 1.4, these devices are typically constructed by infiltrating a light absorbing dye into a nanoporous metal oxide semiconductor matrix, which the dye binds to through an anchoring group, as discussed in Section 1.6. The pores are then filled with an electrolyte redox couple. Under illumination, the dye absorbs light to create excitons, with one carrier type, typically an electron, transferring to the metal oxide while the electrolyte replenishes charges transferred from the dye to the oxide matrix (in other words, accepts the hole generated) and transfers charge to the counter electrode. The charge transfer from the dye to the metal oxide is an ultrafast process with a slow back-transfer, making charge transfer an extremely efficient process.[37] The injection process generates excess carriers in the metal oxide, splitting the Fermi level into electron and hole quasi-Fermi levels (Section 1.3). The splitting between the electron quasi-Fermi level in the metal oxide and the electrolyte Fermi level defines the open circuit voltage (Voc) of the device.[42] Electric fields are effectively screened by the electrolyte solution and carriers are driven to the contacts by diffusion.[42]

The best performing metal oxide matrix is TiO2, and this is the most commonly used material today. However, other metal oxides have been considered. The runner-up is ZnO,

96 which has a higher bulk mobility than TiO2, a similar band alignment with dyes, and the richest family of available nanostructures.[5, 160] Unfortunately, ZnO suﬀers from low charge injection eﬃciencies.[12] One component of this may be the slower charge injection rate into

ZnO, which is 100 times slower than into TiO2.[12, 160] Another significant component may be the formation of Zn-dye complexes that coalesce into aggregates on the surface.[12] The most effective sensitizers bind to ZnO through carboxylic acid groups, and complex formation is a consequence of the same organic acid dissolution process of ZnO discussed in Chp. 4.[9, 70] These complexes are strongly absorbing but are poorly coupled to the surface and do not efficiently inject charge into the ZnO substrate as evidenced by their strong luminescence.[9] Because of the high surface area of the mesoporous metal oxide matrix, just a few multilayers of dye can substantially reduce the incident photon-to-current conversion efficiency (IPCE). The impact of these complexes has been clearly demonstrated by varying the dye concentration in the staining solution.[70] The IPCE peaks at a particular staining time and falls off after this time, and the IPCE can be improved by reducing the dye concentration and increasing the staining time. It is hypothesized that the peak IPCE is due to a trade-off between dye complex accumulation and complete dye coverage of the mesoporous matrix.[70]

The alloy Zn1-xMgxO (ZnMgO) considered in Chp. 4 appears to be a promising alternative to ZnO because of the significant improvement in etch resistance observed under exposure to benzoic acid (BA) even at small (10%) Mg concentration. Ideally this etch resistance would reduce the rate of Zn-dye complex accumulation. The biggest caveat in this hypothesis is the mechanism by which this etch resistance is achieved. As discussed in Chp. 4, the origin of the etch resistance may be the orders-of-magnitude lower lability of the Mg ion compared to the Zn ion, which may increase the etch resistance of the lattice.[72, 158] However, the difference in lability between Zn and Mg may lead to selective etching of Zn that causes the accumulation of a Mg-rich ZnMgO ’scaffold layer’ which limits further diffusion of acid to the subsurface and diffusion of Zn ions out of the scaffold layer. This process

97 may actually be the factor limiting the etch rate. In this mechanism, the etch resistance would only be apparent after the scaffold layer has formed. The initial transient etch rate could be significantly different from the steady-state etch rate observed after the scaffold has formed. However, for the application of ZnMgO in dye sensitized solar cells, it is clear that the initial etch rate is the decisive factor in the applicability of this material due to the sensitivity of performance to just a few layers of metal-dye complex. The etch rate of ZnMgO exposed to BA measured in Chp. 4 is definitively a steady-state etch rate because it involved measuring significant changes to film thickness which would only occur after a scaffold layer has formed. Hence, there is a need to study the effect of Mg on the initial stages of etching and determine whether the steady-state etch resistance carries over to the initial etch rate. The goal of this work is to better understand the initial stages of etching and how the etch resistance may develop over time. To pursue this goal we study the sensitization of the polycrystalline ZnMgO material of Chp. 4 by the ruthenium-based dye cis-bis(isothiocyanato)bis(2,2’-bipyridyl-4,4’- dicarboxylato)-ruthenium(II) (N3 dye, Figure 1.4) which contains four carboxyl anchoring groups. We also perform further characterization of benzoic acid (BA) treated ZnMgO surfaces. The use of N3 dye provides two advantages over the use of the smaller organic acid, BA, studied in Chp. 4. First, we find the etch rate of N3 dye to be significantly lower than that of BA, as will be discussed in the results. Second, the surface concentration can be estimated even at sub-monolayer quantities through simple UV-Vis absorption measurements. These properties make it possible to study the initial etch rate – the etch rate before any scaffold layer has a chance to form – by using the dye aggregate concentration as a proxy for etch rate. This is not possible with BA. The initial etch rate is studied through UV-Vis, PM-IRRAS and photoluminescence (PL) measurements of N3 sensitized ZnMgO. We determine the surface concentration of N3 dye using UV-Vis and correlate this with the results of PL measurements in order to conclude that the majority of the dye molecules on the surface are bound in non-injecting metal-

98 dye complexes. By comparing surface dye complex concentrations across Mg content, we conclude that the initial etch rate actually increases with Mg concentration, and the steady- state etch resistance does not carry over to the initial etching conditions. PL and PM-IRRAS spectra allow us to determine that the complex formed during N3 sensitization of ZnMgO is a Zn-dye complex, with no evidence for the formation of a Mg- dye complex. XPS compositional analysis of BA treated ZnMgO also indicates that Zn concentrations are enhanced over Mg at the surface. This is likely due to the adsorption of Zn- benzoate complexes on the surface. Together, these measurements of carboxylate complexes on ZnMgO indicate that in general Zn complexes form the majority of etch products. As a further test for the formation of a scaffold layer, we studied the sensitization of ZnMgO after it had been pre-treated with a dilute non-bonding mineral acid. The purpose of the pre-etch was to try to form the scaffold layer prior to dye sensitization. The results of sensitizing these pre-etched samples show that the etch rate no longer increases monotonically with Mg content, but instead has a minimum at 5-10% Mg content. The results presented are consistent with a model of ZnMgO etch resistance in which a Mg-rich ZnMgO scaffold layer forms over time, eventually limiting diffusion of etch species to and from the subsurface. The initial etch rate of N3 dye on ZnMgO contrasts with the steady-state etch rate previously measured for BA-treated ZnMgO, indicating that etch resistance is established over time. The differing trends in metal-dye complex concentration on pre-etched and as-deposited ZnMgO films is also consistent with this idea. Finally, the identification of surface complexes as being primarily Zn-based suggests selective leaching of Zn may occur during the etch process as well.

5.2 Experimental Methodology

ZnMgO was deposited onto gold-on-glass (with Ti adhesion layer), silicon, and quartz substrates by the sol-gel process described in Section 2.1 according to the details found in the experimental section of Chp. 4 (Section 4.2). We restricted our consideration of Mg contents to between 0 and 20% because of the unreliable ﬁlm quality of the 30% samples, as discussed

99 in Chp. 4 (Section 4.3). To produce N3 sensitized ﬁlms, the samples were stained in a 0.5 mM solution in ethanol for the desired amount of time, rinsed with ethanol upon removal,

and dried with flowing N2. The solution concentration and solvent are consistent with typical staining solutions found in the literature.[70, 83] We found that dye also adsorbed to the quartz backside, and this dye was removed with a Q-tip soaked in potassium hydroxide (KOH) base solution. For some experiments, a pre-etch was performed on the bare ZnMgO samples. The samples were dipped in a 0.1 mM HCl solution in water for 30 s. UV-Vis measurements were performed before and after the pre-etch in order to verify that the film experienced a small thickness loss due to etching. UV-vis absorption spectroscopy was performed on the untreated and N3 treated ZnMgO films on quartz substrates using a Cary V spectrophotometer. The spectra were taken between 200 and 600 nm and were used to determine the bandgap of the ZnMgO film and the surface concentration of N3 dye on ZnMgO (discussed below). The molar absorptivity as a function of wavelength (ε(λ)) of N3 dye in ethanol was determined from UV-Vis measurements. The absorbance of a 50 µM solution of N3 in ethanol in a cuvette was measured using an empty cuvette as background. The path length of the light beam through the solution in the cuvette was measured with a caliper, and the absorbance was divided by the concentration and the path length to yield ε(λ). The results of this procedure are shown in Figure 5.1. Three peaks are observed in the spectrum. The two lowest energy peaks are attributed to metal-to-ligand charge transfer (MLCT) electronic transitions. According to theoretical modeling, the metal component of the transition is actually associated with a hybrid Ru – NCS (isothiocyanato group) orbital.[85] The highest energy peak is an intraligand π − π∗ transition occurring on the bipyridyl ligands.[85] Our calculated peak molar absorptivities are in agreement with the literature.[161] The surface concentration of N3 on ZnMgO was estimated using the molar absorptivity of the π − π∗ transition peak near 315 nm. This peak was used because it has the strongest signal. The calculation is not exact because it requires that we assume that the dye has the

100 Figure 5.1: Molar absorptivity (ε(λ)) of N3 dye in ethanol solution measured by UV-Vis absorption. The electronic transitions resulting in each peak are labeled.[85] MLCT = Metal- to-ligand charge transfer. same isotropic angular distribution on the surface as it does in solution. While dye molecules bonded to the surface likely have a preferred orientation, the Zn-dye aggregates of interest in this chapter likely have a more isotropic distribution. The chemical environment can also affect the molecular absorption and comparing absorption in solution to absorption on a surface in air also likely introduces a small error, including solvent effects and local field effects due to depolarization which depend on concentration. To calculate the surface concentration, UV-Vis absorbance measurements were taken before sensitization, after sensitization, and after removing the dye from the surface by dipping for 10 s in a 0.01 M KOH solution in water. Removing the dye after sensitization made it possible to determine whether etching of the ZnMgO had occurred and provided a better background for calculation of the surface concentration. We verified that the dip in KOH had negligible impact on the absorption of the ZnMgO film. The absorbance of the dye alone was calculated by subtracting the bare ZnMgO spectra from the sensitized ZnMgO spectra. Calculations were performed using both the pre-sensitization bare ZnMgO and post dye removal spectra as background. To obtain the surface concentration, the peak dye absorbance near 315 nm was divided by the peak molar absorptivity near that wavelength. The surface concentrations calculated using either background method were in agreement for sensitizing times up to a few hours, indicating

101 little influence from the etching process on ZnMgO film thickness. PL measurements of bare and N3 sensitized ZnMgO samples on quartz substrates were taken using a custom built system. The 514.5 nm laser line of an argon ion laser was used to excite the sample, because this line falls near the peak of the first MLCT transition of N3. A 514.5 nm bandpass filter was employed to block other wavelengths generated in the laser. A 570 nm longpass filter was used to prevent laser light from entering the detector. The photoluminescence generated by the laser was dispersed by an Acton Research Corporation spectrograph (model 300i) and detected by a Roper Scientific CCD detector (model 7346- 0005). The spectra were taken over the range of 465 to 1030 nm. The resulting spectra were collected and processed using WinSpec software. The laser power was measured immediately after each measurement and each spectrum was divided by its corresponding laser power to generate power-independent spectra, assuming a linear response of the dye with intensity. At least five spectra at five separate points were taken for each film. Polarization Modulated - InfraRed Reflectance Absorbance Spectroscopy (PM-IRRAS) measurements were performed on untreated and N3 treated ZnMgO films. PM-IRRAS provides an extremely surface-sensitive IR absorbance measurement by minimizing environmental noise and emphasizing surface-localized species through polarization modulation of the incident light.[126, 146] See the discussion in Section 2.6. Measurements were performed on samples on gold-on-glass substrates using a Thermo Scientific Nicolet 6700 FT-IR spectrometer with a Nexus PEM Module. Spectra were taken between 650-4000 cm-1. Each alloy was measured before and after N3 treatment. The spectra of the N3 treated surfaces were divided by spectra from their respective untreated surfaces to generate an absorbance spectrum.[88]

For XPS measurements, ZnO and Zn0.8Mg0.2O samples on Si were prepared as discussed above. In accordance with the procedure of Section 4.2, the samples were exposed to a 2 mM BA solution in hexane for half an hour, and rinsed with hexane upon removal in order to modify the metal oxide surface and produce an etched surface. XPS measurements were

102 taken before and after BA modiﬁcation. The measurements were taken using 1486.7 eV Al

Kα monochromatic x-rays. Full spectrum scans of each sample were recorded, followed by high-resolution scans of the following orbitals: Mg 1s, Zn 2p(3/2), C 1s, O 1s, Mg 2p, and Zn 3d. The binding energies (BE) and sensitivity factors of the Zn, Mg, and O orbitals used in this study are shown in Figure 5.2. The sensitivity factor is the relative signal intensity expected from an orbital compared to a standard orbital, and is used in the calculation of sample composition. It depends in part on the instrument used, including the electron lens. The composition of each sample in terms of Zn, Mg, and O atoms before and after BA treatment was determined from the integrated intensity of each high-resolution orbital scan using the following pairings: (i) Mg 1s, Zn 2p(3/2), and O 1s, and (ii) Mg 2p, Zn 3d, and O 1s. These pairings take advantage of the diﬀerent escape depths (inelastic mean free path) of high and low BE orbitals (as indicated in Figure 5.2) in order to give information about the composition at diﬀerent (near-surface) depths.[94]

Figure 5.2: Sensitivity factors of the orbitals measured during the XPS experiment as a function of their binding energy. Note that with decreasing binding energy, the escape depth of the electrons increases. Therefore, the Mg 1s and Zn 2p orbitals are more surface sensitive than the Zn 3d and Mg 2p orbitals.

103 5.3 Results and Discussion

ZnMgO films of Mg content 0 – 20 %, in steps of 5 %, were produced according to the sol-gel process discussed in Section 4.2. The bandgaps of the films were determined from UV-Vis measurements according to the procedure of Chp. 4, yielding results consistent with that chapter (Figure 5.3). There is some deviation at higher Mg contents from the results of Olson et al. on which the ZnMgO synthesis procedure is based.[18] The modification of bandgap with Mg content indicates that a true alloy is being formed.

Figure 5.3: Bandgaps for samples studied in this chapter (black circles) compared to the results of the previous chapter (green circles) and the results of Olson et al. (blue circles).[18] The error bars indicate the standard deviation from sample to sample.

ZnMgO samples were sensitized with N3 dye by staining in a 0.5 mM solution in ethanol for one hour. This staining time is consistent with what Keis et al. observed to be optimal for creating ZnO DSSCs using a 0.5 mM solution.[70] UV-Vis and PL were performed on these samples as described in the experimental methodology section, and the results of these measurements are shown in Figure 5.4. The UV-Vis spectrum observed matches with the measured molar absorptivity, except that the MLCT peaks show a blue shift compared to

104 the dye in solution. This blue shift could be due to the transition from solution to a surface in contact with air and/or to shifts in electronic structure due to bonding of the carboxyl groups. The PL spectrum matches the ﬂuorescence spectrum observed by Horiuchi et al. in their study of Zn-dye complexes formed during exposure of ZnO to N3 dye, though our peak wavelength is slightly red-shifted which may be due to system response.[9] This is true regardless of Mg content, as can be seen by plotting the normalized PL spectra together (Figure 5.5).

Figure 5.4: UV-Vis and PL spectra of ZnMgO treated with N3 dye for 1 hr, for Mg contents from 0 - 20 mol. %. A scaled molar absorptivity spectrum is included with the UV-Vis measurements.

The surface density of N3 dye molecules on the ZnMgO samples has been estimated from the absorbance curves of Figure 5.4 according to the procedure given in the experimental methodology, and is shown in Figure 5.6. The integrated intensity of the PL spectra in Figure 5.4, which is proportional to the number of non-injecting dye molecules, is also shown in Figure 5.6. The approximate surface coverage expected from a surface-bonded monolayer of dye on the surface (∼ 5 × 1013 cm-2) is indicated by the dotted line labeled ’monolayer’ in Figure 5.6. This density is estimated from the density given in the literature for dye sensitization of TiO2, which does not form metal-dye surface complexes.[161] Before proceeding,

105 Figure 5.5: The PL spectra of Figure 5.4, normalized to a height of 1, for Mg contents from 0 - 20 mol. %.

we note that there may be a background present in our UV-Vis spectra inﬂuencing the surface concentrations calculated in Figure 5.6. This could be due to scattering induced by the accumulation of dye on the surface. To address this concern, we have also calculated surface concentrations for these samples after baselining the peak used for the calculations (and the molar absorptivity peak) in order to remove any potential inﬂuence of a background. This

approach results in the same trends, but smaller overall coverages by about 1/3. The surface density of N3 dye on our samples is more than an order of magnitude higher than the expected monolayer coverage indicating that most of the dye on the surface is bound up in the form of metal-dye complexes. The agreement between trends in surface concentration and PL intensity support this conclusion. However, the surface concentration is comparable to the Zn atomic surface density which is on the order of 1015 cm-2. We conclude from these results that very little etching has taken place, and the measurements made are a reﬂection of the initial etch rate of the ﬁrst few atomic layers of ZnMgO. One potential concern about these measurements is whether the surface dye concentration may

106 Figure 5.6: Surface coverage and PL integrated intensity of ZnMgO treated with N3 for 1 hr, for different Mg contents. The dotted line labeled ’monolayer’ indicates the expected surface coverage at monolayer coverage of N3 dye, based on reports in the literature for the coverage of TiO2, which does not form metal-dye complexes.[161] have saturated before the etch was completed, in which case the surface dye concentrations may not be a good proxy for the etch rate. However, we have observed the progression of surface concentration over time and have found that surface concentrations continue to increase at longer etch times, indicating the surface concentration has not saturated. We contrast these surface concentrations to the etching observed by BA in Chp. 4, where significant film loss was measured for the ZnO and Zn0.9Mg0.1O samples after 1 hr of etching, to conclude that the etch rate of N3 is significantly slower than BA. The most notable feature of Figure 5.6 is the increasing non-injecting dye concentration on the surface with increasing Mg content, as indicated by both direct surface concentration measurement and PL integrated intensity. One of the goals of this work is to determine whether the steady-state etch resistance of ZnMgO carries over to its initial etch rate. It is clear from Figure 5.6 that exactly the opposite is true when sensitizing as-prepared ZnMgO. We conclude from this experiment that the initial etch rates of N3 on ZnMgO follow the

107 opposite trend to that observed for the steady-state etch rates of ZnMgO exposed to BA discussed in Chp. 4. Since ZnMgO surfaces are composed of both Zn and Mg metal atoms, the possibility exists for the formation of at least two different complexes. The identity of the surface complexes observed is now discussed. The consistent location of the PL peak in Figure 5.5 across Mg contents and its agreement with the PL peak observed by Horiuchi et al. for Zn- dye complexes suggests that all Mg percents form Zn-dye complexes as their primary etch product.[9] However, it is possible that Mg-dye complexes could also luminesce at the same peak wavelength. In Chp. 4, we found that PM-IRRAS was very sensitive to the bonding configuration of the carboxyl groups of BA. We have therefore employed PM-IRRAS of N3 treated ZnMgO as a more effective technique for identifying the complexes formed during sensitization. Figure 5.7(a) shows the results of PM-IRRAS measurements of ZnMgO treated in a 0.5 mM N3 solution for 3 hr. We found that PM-IRRAS intensity does not correlate well with surface coverage determined by UV-Vis measurements and therefore only peak locations and feature shapes are considered in our analysis. The carboxyl stretch region of

− − N3 is shown in Figure 5.7 and the symmetric (νsym(CO2 )) and asymmetric (νasym(CO2 )) stretch modes of the carboxylate ion are labeled. These modes have been identiﬁed from a previous infrared study of N3 sensitized ZnO by Keis et al.[70] The locations of these modes and the rest of the modes observed in Figure 5.7 are in agreement with the IR measurements of N3 sensitized ZnO and synthesized Zn-N3 complexes studied by Keis et

al.[70] The νsym and νasym modes of Figure 5.7(a) do not show signiﬁcant changes in feature shape as a function of Mg content. This is in contrast to the signiﬁcant feature changes we observed in Chp. 4 for PM-IRRAS measurements of BA treated ZnMgO as a function of Mg content. We would expect to see similar trends here if both Zn and Mg complexes were present. Based on this and the agreement in peak location with previously reported IR measurements of Zn-N3 complexes, we suggest that Zn-dye complexes are the only product of etching formed throughout the range of Mg content studied. We also note that the spectra

108 in Figure 5.7(a) show no intensity at the location of the free carboxylic acid stretch mode (ν(COOH), labeled in Figure 5.7(b)). This suggests that all four carboxyl groups in the N3 molecule are bonded. The free carboxylic acid stretch mode is observed in samples sensitized for 24 hr (Figure 5.7(b)), likely due to trapping of unbonded dye by the large Zn-dye complex agglomerates that form during such long sensitizing times.[9]

Figure 5.7: PM-IRRAS Spectra of ZnMgO exposed to N3 dye for (a) 3 hr and (b) 24 hr. The region shown is the carboxyl stretch region of the binding groups on the dye. The vertical − − lines mark the symmetric (νsym(CO2 )) and asymmetric (νasym(CO2 )) stretch modes of the bonded carboxyl groups. The carboxyl stretch of the free acid (ν(COOH)) is also marked in (b).

To make our conclusions about the formation of surface complexes more general, we revisited the modiﬁcation of ZnMgO with benzoic acid (BA). In Chp. 4, we observed the presence of Zn-benzoate etch products in our IR measurements, but did not identify any Mg- benzoate complexes. If Zn-benzoates are the dominant etch product, then the surface should

109 show enhanced Zn concentration compared to Mg after treatment. X-ray photoelectron spectroscopy (XPS) is a highly surface selective tool, sensitive to material within ∼10 nm of the surface.[94] It is also capable of determining elemental surface composition. We employed XPS to study the surface composition of ZnMgO before and after exposure to BA. Samples

of ZnO and Zn0.8Mg0.2O were measured by XPS before and after exposure to 2 mM BA in hexane for 30 min., in congruence with the surface modiﬁcation procedure of Chp. 4. High-resolution spectra of the Mg 1s, Zn 2p(3/2), Mg 2p, Zn 3d, and O 1s orbitals were collected as described in the experimental methodology section. The integrated intensities of the orbital features produced were used to determine surface composition. The escape depth of an ejected core electron, or the depth from which a free electron can escape the material without inelastic scattering, increases as its kinetic energy (KE) increases.[94] Since XPS uses monochromatic x-rays the electron KE decreases as binding energy (BE) increases. Therefore, higher BE orbitals have shallower escape depths, and vice versa for lower BE orbitals. While XPS is already very surface sensitive, the varying escape depths can be exploited to get information about surface composition as a function of near-surface depth. For reference, the escape depth ranges from a minimum of about 2-3 nm for photoelectrons of KE 20 eV to about 20-40 nm for electrons of KE 1000 eV.[94] For this reason, the high BE orbitals Mg 1s and Zn 2p(3/2) provide a more surface sensitive signal than the low BE orbitals (Mg 2p and Zn 3d). Note that only one orbital (O 1s) is used to study oxygen and it is employed in both the low BE and high BE analyses.

Measurements of ZnO and Zn0.8Mg0.2O before and after BA treatment (Figure 5.8(a) and (b)) show that the surface oxygen composition increases relative to the metals after treatment. We attribute this relative increase in oxygen to the carboxyl groups of BA that are present on the surface after treatment as both surface bonded species and Zn-benzoate complexes. The diﬀerence in escape depth between the O 1s orbital and the metal orbitals is not accounted for in the calculation of surface composition and this explains some of the variation observed in oxygen content between the high and low BE analyses.

110 Figure 5.8: XPS composition measurements of ZnO (a) and Zn0.8Mg0.2O (b) before and after treatment with 2 mM benzoic acid (BA) for 30 min. Compositions were calculated using both the high and low binding energy (BE) Zn and Mg orbitals. Note that the O 1s orbital is common to both analyses.

111 The relative Zn and Mg concentrations of Zn0.8Mg0.2O were also compared (Figure 5.9) before and after BA treatment. These compositions were calculated by renormalizing the compositions from Figure 5.8 for Zn and Mg only. The lower bonding energy orbitals show no change in the relative Zn and Mg compositions. An enhancement of the Zn concentration compared to Mg is seen in the high BE orbitals but not in the low BE orbitals, suggesting that Zn is concentrated at the surface. This indicates that the Zn benzoate complexes observed in the IR spectra of Chp. 4 are the dominant complex species on the ZnMgO surface.

Figure 5.9: XPS composition measurements of Zn and Mg before and after treatment of Zn0.8Mg0.2O by BA for 30 min., as described in the procedures.

The PL, PM-IRRAS, and XPS measurements of N3 and BA treated ZnMgO in this chapter and Chp. 4 all point to the idea that only one type of etch product is present on the surface: Zn-carboxylate complexes, with no evidence for Mg based complexes. Since the initial etch rate was observed to be opposite the observed steady-state etch rate, the question of whether the surface conditions of the steady-state etch rate can be engineered through an acid pre-etch of the ZnMgO ﬁlms is explored. ZnMgO samples were pre-etched by dipping in 0.1 mM HCl solution in water for 30 s. This produces an observable change in absorption, but does not signiﬁcantly reduce the thickness of the sample. The samples were then consecutively sensitized in a 0.5 mM N3 ethanol solution for 5 min, 25 min, 30 min,

112 and 1 hr in order to monitor the progression of the etching process. The surface coverage was then determined from UV-Vis measurements (shown for 1 hr exposure in Figure 5.10). The results of this experiment are shown in Figure 5.11. The coverage vs. Mg content curves show a clear departure from the shape of the as-prepared samples of Figure 5.6. In this case, a minimum surface coverage is observed at 5 - 10% Mg content. While these results do not indicate improvements in etch resistance, they do suggest that the pre-etch has substantial inﬂuence on the sensitization process.

Figure 5.10: UV-Vis absorbance spectra of pre-etched ZnMgO samples exposed to a 0.5 mM solution of N3 dye in ethanol for 1 hr, for Mg contents from 0 to 20 %.

Further insight can be found by comparison to samples that have been sensitized for a very long time. Figure 5.12 shows the surface coverage of ZnMgO samples sensitized for 2 hr after being pre-etched compared to the surface coverage of a sample with no pre-etch sensitized for 24 hr. The two curves show the same minimum, but the 24 hr sensitization minimum is more dramatic. This implies that the pre-etch has an eﬀect similar to a long exposure to N3 dye. The diﬀerence between the two curves can be explained as follows: the two treatments both achieve the same trend in dye complex concentration through acid etching of the surface. However, the pre-etch accomplishes this without exposing the surface

113 Figure 5.11: Surface coverage as a function of total exposure time to 0.5 mM N3 dye for different Mg content pre-etched ZnMgO samples. The pre-etch was performed by exposing the films to a dilute mineral acid. The solid lines connect surface coverages calculated using the as-deposited ZnMgO films as background. The dotted lines connect surface coverages calculated using as background bare ZnMgO films that have had their dye removed by submersion in a basic solution.

114 to dye. Subsequent sensitization results in a lower surface dye concentration but the same trend. The 24 hr N3 exposure also etches the surface, but over a much longer time and with subsequent formation of Zn-dye complexes resulting in substantially higher concentrations of dye that may be saturation-limited.

Figure 5.12: N3 surface coverage of a ZnMgO sample exposed to 0.5mM N3 dye for 24 hr and a sample ﬁrst etched with mineral acid and then exposed to N3 for 2 hr. Both show minima at 10% Mg.

Minima generally imply trade-offs between variables, and the minima of Figure 5.11 and Figure 5.12 can be interpreted in this way. A qualitative model based on our working hypothesis that a Mg-rich ZnMgO scaffold formed during etching limits the etch rate is employed. The model is as follows: At 0 % Mg content, the etch rate during dye sensitization is always high regardless of a pre-etch or lack thereof, leading to significant Zn-dye complex accumulation. However, the addition of Mg allows a scaffold layer to form during etching (whether by mineral acid or by dye), leading to improved etch resistance to N3 and reduction of aggregate formation. The results of Chp. 4 are assumed so that the higher Mg content films afford higher etch resistance and therefore lower Zn-dye complex formation. In order for this to hold, as Mg content increases, the scaffold layer must become either denser or

115 thicker to further slow diffusion to the subsurface. The dye molecules may attach to this scaffold layer as well. Higher Mg content scaffold layers will accommodate more dye because of their increased surface area. There is therefore a trade-off between etch resistance, which decreases dye complex accumulation as Mg content increases, and attachment of dye to the scaffold, which increases as Mg content increases. At high Mg contents, the attachment of dye to the scaffold layer outcompetes the reduction in Zn-dye complex formation, leading to higher surface coverages.

5.4 Conclusions

Surface modification of ZnMgO films by N3 dye and benzoic acid (BA) has been employed as a vehicle for pursuing the goal of understanding better the initial stages of etching in this material and how the observed steady-state etch resistance may develop over time. Such an investigation informs the applicability of ZnMgO as a replacement for ZnO in applications such as dye sensitized solar cells where etch resistance is desired and initial etch rates are the critical rate. In Chp. 4, we proposed a mechanism of etch resistance in which Zn is selectively leached from the surface, leaving behind a Mg-rich ZnMgO scaffold that limits the etch rate by limiting diffusion of etch species to and from the ZnMgO subsurface. This work further informs that idea. The relatively slow etch rate of N3 dye, combined with its strong optical absorption made it possible to study the initial, transient etch rate of the ZnMgO film, in contrast to the steady-state etch rates measured in Chp. 4. The UV-Vis, PL, PM-IRRAS, and XPS measurements of as-prepared ZnMgO sensitized with N3 dye or modified with BA paint a distinct picture of the initial etching process. During etching of the first few atomic layers, the etch rate of ZnMgO actually increases with Mg content, the opposite of the steady- state trend, indicating that an initial etch period is required to establish the steady-state etch resistance observed in Chp. 4. Photoluminescence and PM-IRRAS measurements of the sensitized ZnMgO surface show that the etch products produced are primarily Zn- dye aggregates, regardless of Mg content. There is no evidence suggesting that Mg-dye

116 aggregates form as well. A similar conclusion can be made about the products of etching by benzoic acid. XPS showed that the concentration of Zn on the surface is enhanced relative to Mg after exposure of ZnMgO to benzoic acid, suggesting that Zn benzoates, as opposed to Mg benzoates, are also the primary product of etching by benzoic acid. The formation of primarily Zn-based etch products using two different modifiers suggests the general conclusion that the etch products of carboxylic acids on ZnMgO will be primarily Zn-carboxylates. The observation of primarily Zn-based etch products also implies that either Mg is transferred directly to solution or that the Mg is still present in the surface. These results are therefore consistent with the idea that Zn is selectively leached from the film as etching progresses. Taken together, these results are consistent with an etch resistance mechanism in which leaching of Zn from the film over an initial etch period establishes a Mg-rich ZnMgO scaffold that limits the steady-state etch rate. Given that as-deposited ZnMgO does not show a reduction in Zn-dye complex formation compared to ZnO, the dilute mineral acid pre-etch of ZnMgO suggests a direction for future development of this material for organic acid surface modification applications such as dye sensitized solar cells. The surface coverage of N3 on pre-etched ZnMgO films showed a minimum at 10% Mg content instead of the monotonic increase in N3 concentration observed on the as-deposited films. The appearance of this minimum indicates that a reduction in dye complex aggregate formation and initial etch rate may be achievable by optimizing a pre-etch process. The more pronounced minimum observed for ZnMgO exposed to N3 for 24 hrs suggests that a more dilute or weaker acid may be more effective at establishing an effective leached layer.

5.5 Acknowledgements

T.M.B. and E.P.M. would like to thank Gaute Otnes and Chito Kendrick for training on photoluminescence measurements and help troubleshooting. This material is based upon work supported by the National Science Foundation through Grant No. DMR-0907409 and through the Renewable Energy Materials Research Science and Engineering Center under

117 Grant No. DMR-0820518. DCO’s funding is provided by the U.S. Department of Energy under Contract No. DE-AC36-08-GO28308 with the National Renewable Energy Laboratory through the DOE SETP program.

118 CHAPTER 6 CONCLUSIONS

Organic acid dissolution of ZnO is an important problem in dye sensitized and organic solar cells in which organic acid surface modifiers are commonly used to tune surface properties. The etching process reduces the quality of monolayers formed on ZnO by leaving behind etch products and damaging the ZnO layer itself. Two ways to address this issue have been explored in this thesis. The first was to employ a covalent attachment scheme instead of an acid/base reaction for attaching modifier molecules to the surface. The second was to consider an alloy of ZnO that is more resistant to etching but preserves the electronic properties of ZnO.

6.1 Project Conclusions

The covalent bonding scheme explored is that of the triethoxysilane (TES) attachment. Chp. 3 culminates several years of experience within this research group of attaching TES- based molecules to ZnO. Our work with the TES attachment scheme has demonstrated that it is less reactive than other silanes, and an effective technique for modifying ZnO.[55, 69, 78] Going a step further, Chp. 3 demonstrates definitively that the TES attachment scheme can be used to effectively tune metal/oxide organic interfaces as well as photovoltaic device properties. We demonstrated a work function tunability of 600 meV and a Voc tunability of 95 mV when ZnO was used as the electron collecting contact in an inverted bulk heterojunction (IBHJ) OPV device. The study also raised a number of interesting questions concerning the microscopic properties of mixed monolayers as well as the relationship between the contact work function and the Voc of BHJ devices. We observed a 0.14 V/eV relationship between the contact work function and the Voc of our devices. This raises the question of why this value is observed. A number of effects may be contributing to this value in our devices. First of all, there

119 may be Fermi level pinning effects where the metal oxide Fermi level becomes pinned to the LUMO states of the organic.[57, 59] Other potentially important possibilities arise from the fact that we used mixed monolayers to tune the metal oxide/organic interface. There may be microscopic effects having to do with pinholes in the monolayer or other defects that significantly impact the electronic properties of the surface.[127, 135, 136] Furthermore, the use of mixed monolayers may produce segregated domains of differing Fermi level, as has been observed in other mixed populations of adsorbed species, and this may also have an impact on the electronic properties of the interface.[128] All these questions warrant further investigation and will be discussed further in Section 6.2. Another important conclusion we’ve come to about the use of TES monolayers is that it is relatively difficult to obtain TES molecules which lower the work function of ZnO, a necessary attribute for employing these molecules to improve the contact properties of ZnO in typical electron collecting OPV applications. Moreover, even for molecules like PTES that do lower the work function, once used within a PV device they still produce a lower

Voc suggesting that a detrimental dipole forms upon interface formation. Our previous work comparing octadecyltriethoxysilane to octadecanethiol in hybrid ZnO/P3HT bilayer devices also suggests that the TES modification process induces a work function-raising dipole.[55] Therefore, TES modifiers may find more effective application in hole-collecting metal oxides such as NiO or MoO3 where larger work functions are desirable for improving Voc. The technique of forming etch-resistant alloys of ZnO is another promising avenue for reducing the impacts of etching on the metal oxide without creating fundamental changes in the electronic properties. In Chp. 4 we observed decreasing steady-state, macroscopic etch rates of wurtzite Zn1-xMgxO (ZnMgO) alloys exposed to benzoic acid as Mg content increased. An order of magnitude reduction in etch rate was observed at 20% Mg content and above. We developed a paradigm for the origin of the etch resistance in ZnMgO which also allowed us to identify other possible alloys of potentially higher etch resistance. We suggested that the orders-of-magnitude lower lability of the Mg ion compared to the Zn ion

120 may provide lower reactivity to the lattice and reduce the ability of acids to transfer ions to solution. As a consequence of such a model, even lower lability ions with the same valence as Zn may be considered for the creation of alloys. This leads to the prediction that wurtzite

Zn1-xNixO (ZnNiO) could have even higher etch resistance than ZnMgO. Possible routes for studying ZnNiO are discussed in Section 6.2. However, this simplistic picture of etch resistance in which the reactivity of the lattice as a whole is reduced has been brought into question through the deeper investigation of the mechanism in Chp. 5, as discussed below. One of the potential applications of the etch resistance of ZnMgO is to reduce dye aggregate formation in ZnO dye sensitized solar cells (DSSCs) by replacing the ZnO with ZnMgO. This could boost the incident photon-to-current conversion efficiency. While the lability of Mg may be the underlying source of the improved etch resistance, the mechanism by which Mg addition imparts this etch resistance is an important process to understand if this property is to be employed practically. It is quite common to find that etching is selective in mixed metal oxides, leaching out the more reactive metal ions and leaving behind a leached layer of the less reactive oxide.[71, 72] This process can actually limit the etch rate by limiting the diffusion rate of molecules toward and away from the mixed metal oxide surface below the leached layer. We explored this hypothetical mechanism of etch resistance in Chp. 5, in the context of dye sensitization of ZnMgO. The initial etch rate in ZnO based DSSCs is quite important because a multilayer just a few molecules thick can significantly impact the performance of the device. However, in our model of leach-layer limited etching, the initial etch rate could be quite different from the steady state, macroscopic etch rate. We found that the initial etch rate of N3 dye on ZnMgO increased with Mg content, in contrast to the steady-state etch rate of benzoic acid on ZnMgO measured in Chp. 4. Our identification of surface metal-dye complexes as Zn-dye complexes, regardless of Mg content, combined with XPS measurements suggesting Zn-benzoates are the primary etch product of benzoic acid- treated ZnMgO, suggest that Zn is selectively leached from the surface. We conclude from these observations that the etch resistance of ZnMgO likely arises from a surface condition

121 that develops over time. Furthermore, our results are consistent with a leached-layer model of etch resistance. Based on the observation that the initial etch rate of N3 on ZnMgO differed from the steady-state etch rate of BA-treated ZnMgO, we tried pre-etching the ZnMgO surface before sensitization by exposure to a dilute mineral acid which does not adsorb to the surface. We discovered that the surface coverage no longer increased monotonically with Mg content, but showed a minimum at 5-10% Mg. While this does not demonstrate improved etch resistance, the observed minimum is a promising step forward. We can interpret the results within the leached layer model: the minimum is due to a trade-off between improved etch resistance due to a leached layer that increases in depth or density with Mg content and attachment of dye to the leached layer, increasing the surface concentration. These results raise some important questions about this pre-etch technique. Could the pre-etch be optimized by using a more dilute acid or even an organic acid? Could organic acids prove more effective in forming a leached layer because of their tendency to isolate surface atoms through complex formation? Could an adsorbing acid be used in pre-etching, and then still allow dye to subsequently attach to the surface? Finally, how would a leached layer affect interface properties in organic or dye sensitized solar cells?

6.2 Future Project Suggestions

Understanding the influence of the contacts on the Voc of bulk heterojunction devices is of substantial interest to the organic photovoltaics community. Chp. 3 suggests that monolayer modification is an effective technique for tuning contacts and understanding the behavior of these contacts under changing energy level alignment. Opportunities for further pursuit of this tactic exist in molecules with stronger dipoles and/or higher surface densities, since the increased tunability will provide a stronger signal of the relationship between the metal oxide work function and the interfacial energy level alignment. Chp. 4 suggests that a monolayer of carboxylic acid is formed on ZnMgO prior to etching and remains on the surface during etching, and previous work has demonstrated the effectiveness of carboxylic acids in

122 introducing dipoles on the ZnO surface.[99, 121] Therefore carboxylic acids are likely an effective attachment group for introducing large surface dipoles, and can be employed with the results of Chps. 4 & 5 in mind. A study of mixed monolayers of dipolar carboxylic acids modifying ZnO (or perhaps another metal oxide) in bulk heterojunction OPV devices may produce stronger conclusions about how the metal oxide work function affects device properties. However, the concerns raised in Chp. 3 about defects, inhomogeneity, or other microscopic properties of single component or mixed monolayers need to be considered. One route to addressing these concerns may be found in scanning Kelvin probe force microscopy (SKPM) measurements. While Kelvin probe (Section 2.7) can only measure spatially aver- aged surface potentials, modern atomic force microscopes can be used to take nanoscale surface potential measurements through SKPM. Previous work with this technique has shown phase segregation and surface potential contrast of molecular monolayers at the nanoscale.[128] It would likely be worthwhile to investigate the use of this technique for studying potentially inhomogenous or mixed monolayers, with the ultimate goal in mind of matching microscopic properties to macroscopic interface behavior. One of the major prospects from Chp. 4 was the prediction that ZnNiO may be more resistant to etching than ZnMgO. This material has been synthesized previously through solution methods, and its high bandgap semiconductor properties established.[75, 159] This material may prove to have some advantages over ZnMgO beyond its potentially higher etch resistance. For instance, it may not display the same exponential increase in resistivity as ZnMgO.[17, 18] It should be possible to establish a simple sol-gel synthesis route of ZnNiO. Other researchers have produced ZnNiO films using acetates of both Zn and Ni.[75, 159] A process using nickel formate, which has been employed to form NiO films, and zinc acetate, which has been used to form ZnO in this thesis, might also be considered.[63] The study of the synthesis and properties of ZnNiO for use as an etch resistant metal oxide semiconductor in organic electronics could potentially be a fruitful long term project. However, understanding

123 its etching mechanism would again be of significant importance for its applicability in organic electronic devices. The projects in this thesis have provided an understanding of both covalent (TES) and acid/base (carboxylic acid) attachment schemes. Such an understanding makes it possible to begin studies comparing these types of bonding methods. Our previous work has speculated about the dipoles of the surface bond of different attachment schemes.[55] The fact that TES molecules bond covalently through an oxygen atom while carboxyl groups coordinate to Zn atoms may make a comparative study of surface dipole and interface charge transfer effects using TES another potentially fruitful project.

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139 APPENDIX A - FURTHER EXPLANATION OF METHODOLOGY

A.1 Infrared Active and Inactive Modes: An Example

One of the simplest examples illustrating the requirement that the dipole moment of a chemical structure change during an infrared-active vibration are the symmetric and asym-

metric stretch modes of the carbon-oxygen bonds in carbon dioxide, CO2 (Figure A.1).

Overall, CO2 has zero net dipole moment, though charge is shared unequally between the oxygen and carbon atoms. In the symmetric stretch of CO2, the synchronized stretching and contraction of the C-O bond leaves the net dipole at zero throughout the vibration cycle making the mode infrared inactive. However, in the asymmetric stretch the movements of the oxygen atoms are exactly opposite, and the unequal C-O bond lengths results in a net dipole. The asymmetric stretch is infrared active because the net dipole changes throughout the vibration.

Figure A.1: Examples of infrared inactive and active modes of carbon dioxide.

140 A.2 Fourier Transform Infrared Spectrometer Design

FTIR spectrometers have several advantages over dispersive spectrometers. They provide faster spectrum acquisition, making it possible to average many spectra and thus reduce noise. This can be done in less time than it would take to acquire one scan in a dispersive instrument. FTIR spectrometers also have less signal loss during measurement, improving signal to noise ratios. Finally, wavelengths are referenced to an internal laser of known wavelength (and very small linewidth) so that the accuracy of the spectrum is always ensured. FTIR spectrometers use a broadband light source. The light from this source is passed through an interferometer (Figure A.2(a)) that generates a time-dependent interference signal called an interferogram. A moving mirror causes different wavelengths to interfere at different times, coding frequency information into the signal intensity as function of time. For a transmission experiment, Figure A.2(b), the light strikes the sample at normal incidence. The sample absorbs some of the infrared light altering the interferogram. The signal intensity is then measured over time by a detector. The detector signal is processed by a combination of signal electronics and a computer (Figure A.2(c)). The computer performs a fast Fourier transform of the interferogram from the time domain to the frequency domain to obtain the IR spectrum. For PM-IRRAS measurements (Figure A.2(d)), the interferogram is generated in essentially the same way, but the sample measurement process is modified. The advantages of this setup are discussed in Section 2.6. The IR light is first passed through a photoelastic modulator that modulates the polarization of the light. Then the light is reflected off a sample deposited on a conductive substrate. Absorption occurs as the light passes through the sample on reflection. Finally, the light intensity is measured by the detector. The signal is processed differently in PM-IRRAS. A demodulator separates the two polarizations of light into two separate signals based on the frequency of the demodulator. These signals are then Fourier transformed separately and used to generate an absorbance spectrum as discussed in Section 2.6.

141 Figure A.2: Experimental setup of an FTIR spectrometer. (a) Broadband interferometer. (b) Sample and detector geometry for a transmission experiment. (c) Electronics and computer for signal processing. (d) Required components and geometry for a PM-IRRAS experiment.

A.3 Simpliﬁed Schematic of the Kelvin Probe

In order to determine the contact potential difference between two samples, a Kelvin probe can be used as discussed in Section 2.7. The Kelvin probe measures the work function difference between the probe and the sample by applying a voltage (Vappl) and determining at what value of Vappl the potential between tip and sample reaches zero. In order to do this, the Kelvin probe circuit is set up as in Figure A.3. Note that the probe tip and sample form a capacitor, storing charge due to the contact potential difference between the two at a potential V . If the tip is oscillated above the sample, the capacitance will change, and thus the stored charge will be modulated. This generates a current that flows through the current- to-voltage converting resistor RI/V . This current is measured as a voltage drop across the resistor Vout. If Vappl is tuned to the contact potential difference so that V = 0, no charge will be stored and Vout = 0 as well. Thus, the contact potential difference is determined by Vappl

142 when Vout = 0. This is shown in the boxed equations of Figure A.3. Measuring zero voltages is challenging and subject to the noise in the measurement near zero signal. In order to more accurately measure the contact potential difference, Vappl is swept over ∼10 V centered on 0 V and the peak-to-peak voltage of Vout is measured simultaneously. The intercept at which Vout = 0 is then determined from the line produced, and the corresponding Vappl gives the contact potential difference. The off-zero measurements used to determine the contact potential difference in this technique substantially increases accuracy and precision.

Figure A.3: Schematic of Kelvin probe circuit showing how Kelvin probe measurements can be made. A number of other elements for amplifying the signal, etc. are present but are not shown here.

143 A.4 TM-AFM Feedback Loop Schematic

In tapping mode AFM, the tip is oscillated above the sample surface and changes in the amplitude of the oscillation are monitored. Adjustments to the height of the tip are made in order to keep the tip oscillation at a constant amplitude set point. This is accomplished through a feedback loop employing a proportional-integral (PI) controller. The tip oscillation amplitude is monitored through a laser which reflects off the surface of the tip onto a photodiode. The signal on the photodiode is read by the controller, and this signal determines the output action of the controller. If the amplitude is off its setpoint, the controller can adjust the height of the tip above the surface using a sample stage whose height is controlled by a piezoelectric actuator. In some AFM systems, the tip height is adjusted by a piezo that moves the tip itself, instead of the sample. The movement of the sample stage will adjust the tip’s proximity to the surface thereby adjusting the tip’s oscillation amplitude. Changes in topography are thus monitored and can be reacted to through the photodiode-controller-piezo feedback loop (Figure A.4).

Figure A.4: Diagram demonstrating the feedback loop that controls the AFM microscope.

144 APPENDIX B - SUPPORTING INFORMATION FOR CHAPTER 3

B.1 Calculation of Surface Proportion of 4CPTES and PTES from Infrared Spectrum

The first step in the calculation of the proportion of 4CPTES and PTES on a mixed monolayer surface is to identify integrated intensities from the infrared spectrum which are unique to each modifier (that is, there are no contributions from the other molecule). Figure B.1 shows how the two peaks identified in Section 3.3 can be analyzed to find such integrated intensities. It is assumed that the ratio of A and B in the 4CPTES only sample is characteristic of those modes in 4CPTES and is therefore a constant across the sample set. The contribution (B0) of 4CPTES to the peak Γ in mixed samples is then calculated as:

B B0 B = → A0 = B0 (B.1) A A0 A B0 can then be subtracted from Γ to get the contribution (C0) of PTES to Γ: B C0 = Γ − A0 (B.2) A Now that characteristic integrated intensities have been obtained for both 4CPTES and PTES they need to be normalized to account for any diﬀerences in oscillator strength, bonding orientation, etc. that might contribute to the intrinsic absorptivity of the mode. To begin the process, it is noted that the integrated intensities of the modes A0 and C0 are proportional to the number of 4CPTES (n4C ) and PTES (nP ) molecules on the surface, respectively, through proportionality constants α and γ:

0 A = αn4C (B.3)

0 C = γnP (B.4)

It is then assumed for the purposes of this calculation that for the 4CPTES and PTES only

samples the number of molecules on the surface is the same (ntot):

145 Figure B.1: Diagram demonstrating calculation of integrated intensity contributed by PTES to the peak at 1130 cm-1 (boxed, labeled Γ). The peak at 1090 cm-1 is contributed only by 4CPTES and is not present in PTES only samples, while the peak at 1130 cm-1 has contributions from both molecules. Using the 4CPTES-only treated sample, the ratio of the two peaks A and B is taken as constant. Using this constant, the contribution (B0) from 4CPTES to the peak Γ in mixed 4CPTES/PTES samples is calculated from the integrated intensity of the peak A0 which is contributed by 4CPTES only. B0 is subtracted from Γ to ﬁnd the contribution C0 from PTES only. In this way the integrated intensities unique to each modiﬁer (A0: 4CPTES, C0: PTES) can be obtained.

146 A = αntot (B.5)

C = γntot (B.6)

By dividing A0 by A and C0 by C, a quantity independent of oscillator strength, etc. can be obtained:

A0 n = 4C (B.7) A ntot C0 n = P (B.8) C ntot If A0/A and C0/C are added, the following is obtained:

A0 C0 n + n + = 4C P (B.9) A C ntot n4C + nP is the total number of molecules on the sample surface and ideally n4C + nP = ntot. However, experimentally this is not the case and there is variation in the total number of molecules on each sample. This is true of the 4CPTES and PTES samples, but the assumption that they have the same number of molecules allows a convenient and robust mathematical formulation for calculating the fraction of each molecule on the surface. To calculate the fraction of each molecule on the surface, simply divide the normalized intensity due to 4CPTES or PTES by the total normalized intensity:

0 A /A n4C 0 0 = (B.10) A /A + C /C n4C + nP

0 C /C nP 0 0 = (B.11) A /A + C /C n4C + nP For each experimental variable (A, A0, B, and Γ) needed for the calculation, the average value from all experiments was used. The uncertainty in each variable was taken to be the largest deviation from the average across all the experiments. The uncertainty in C’, C and each of the fractions n4C /n4C + nP and nP /n4C + nP was determined from standard propagation of uncertainty.

147 B.2 Water Contact Angle Measurements

Contact angle measurements of the mixed monolayer treated and control ZnO surfaces are given in Table B.1.

Table B.1: Water contact angle (CA) measurements of each of the four surface treatments and the deposition control sample (see Section 3.2 for explanation of this control). The listed contact angle measurements are an average from all experiments. The uncertainty (∆CA) in the measurements is the mean square average of the uncertainties from each experiment. The uncertainty in each experiment was taken as the standard deviation of multiple measurements of the sample. *Average value for CA of UV-Ozone cleaned sol-gel ZnO surface from previous work.[69]

Surface Treatment CA (deg.) ∆CA (deg.) PTES 50.8 2.4 2:1 P:4C 47.9 2.1 1:2 P:4C 45.6 1.7 4CPTES 42.4 3.3 Control 31.1 1.0 UV-Ozone* 17.6 3.6

B.3 Atomic Force Microscopy Measurements

Surface roughnesses of mixed monolayer treated and control ZnO surfaces are given in Table B.2. AFM phase images of these surfaces are given in Figure B.2.

Table B.2: RMS roughness determined from tapping mode AFM topographic scans of treated ZnO surfaces. Reported values are an average calculated from all scans taken and the uncertainties reported are the standard deviation of the roughnesses.

Surface Treatment RMS Roughness (nm) Control 1.6±0.2 PTES 1.2±0.1 2:1 P:4C 2.0±0.1 1:2 P:4C 1.7±0.2 4CPTES 2.0±0.8

148 (a) PTES (b) 2:1 P:4C

(e) Control

Figure B.2: AFM phase images of monolayer treated ZnO surfaces.

149 B.4 Dark J-V Curves

The J-V curves measured in the dark for the solar cells discussed in Section 3.3 and the shunt resistance calculated from these curves at V = 0 are given in Figure B.3 and Table B.3.

Figure B.3: J-V curves of the monolayer treated ZnO P3HT:PCBM devices in the dark.

B.5 Eﬀect of Light Soaking

The eﬀect of light soaking on the J-V curves of our monolayer treated ZnO P3HT:PCBM bulk heterojunction devices is illustrated by 4CPTES in Figure B.4.

150 Table B.3: Shunt resistance (RSH ) calculated from the dark J-V curves (Figure B.3) at V = 0. While there is significant variation in the shunt resistances between treatments, there is no clear trend suggesting that Voc is correlated with shunt resistance in the dark. This is in agreement with Ratcliff et al. [52], who also found no correlation between dark J-V curve shunt resistance and Voc. There is, however, a correlation between the vacuum level shift (contact potential difference) caused by the interface dipole and the diode turn-on voltage as would be expected.

2 Surface Treatment RSH (kΩ cm ) Standard 2.00 Control 0.89 PTES 2.10 2:1 P:4C 0.53 1:2 P:4C 2.40 4CPTES 1.62

Figure B.4: J-V curves of 4CPTES-treated ZnO devices before and after light soaking. A double-diode or s-kink is present before light soaking and is minimized after 20 min. of light soaking. With light soaking, the Voc is correspondingly increased.

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