Novel Uses of Dioxide for Solar Cells

A thesis submitted as partial fulfillment of the requirement for the Degree of

Doctor of Philosophy

by

Bryce Sydney Richards

at the

Centre for Photovoltaic Engineering and the School of Electrical Engineering University of New South Wales Sydney 2052 New South Wales

Australia

April 2002

CENTRE FOR PHOTOVOLTAIC ENGINEERING UNSW Certificate of Originality

I hereby declare that this submission is my own work and that, to the best of my knowledge and belief, it contains no material previously published or written by any other person nor material which to a substantial extent has been accepted for the award of any other degree or diploma of a university or other institute of higher learning, except where due acknowledgment is made in the text.

I also declare that the intellectual content of this thesis is the product of my own work, even though I may have received assistance from others on style, presentation and language expression.

Bryce Richards

Richards, Bryce Sydney Novel Uses of for Silicon Solar Cells PhD Thesis Centre for Photovoltaic Engineering The University of New South Wales Sydney, NSW 2052, Australia Copyright c 2002 all rights reserved

ISBN 0 7334 1971 2 Abstract

Titanium dioxide (TiO2) thin films have a long history in silicon photovoltaics (PV) as antireflection (AR) due to their excellent optical properties and low deposition cost. This work explores several novel areas where TiO2 thin films could be use to enhance silicon (Si) performance while reducing device fabrication costs.

Amorphous, and TiO2 thin films are deposited using ultrasonic spray- deposition (USD) and chemical vapour deposition (CVD) systems, both designed and con- structed by the author. Initial experiments confirmed that no degradation in the bulk minority carrier lifetime (τbulk) occurred during high-temperature processing, although the stability of the USD-deposited TiO2 films was dependent on the furnace ambient.

A major disadvantage of TiO2 AR coatings is that they afford little surface . In this work, a novel method of achieving excellent surface passivation on TiO2-coated silicon wafers is presented. This involved growing a 6 nm-thick SiO2 layer at the TiO2:Si interface by oxidising the wafer after TiO2 film deposition. The increase in surface passivation afforded by the interfacial SiO2 layer results in a decrease in the emitter dark saturation current −14 2 (J0e) by nearly two orders of magnitude to 4.7 − 7.7 × 10 A/cm . This demonstrates the compatibility of the TiO2/SiO2 stack with high-efficiency solar cells designs.

By varying the film deposition and annealing conditions, TiO2 refractive indices in the rangeof1.726 − 2.633 (at λ = 600 nm) could be achieved. Subsequently, a double-layer antireflection (DLAR) was designed comprised of low and high TiO2 material. The best experimental weighted average reflectance (Rw) achieved was 6.5% on a planar silicon wafer in air. TiO2 DLAR coatings are ideally suited to multicrystalline silicon (mc-Si) wafers, which do not respond well to chemical texturing.

Modelling performed for a and ethyl vinyl acetate (EVA) encapsulated buried-contact solar cell indicated that a TiO2 DLAR coating afforded a 7% increase in the short circuit current density, when compared to a standard, commercially-deposited TiO2 single-layer AR coating.

Finally, it is demonstrated that chemical reactions with prevent TiO2 from acting as a successful phosphorus diffusion barrier or dopant source. The applicability of TiO2 thin films to various silicon solar cell structures is discussed. Acknowledgements

Many people contributed to the success of this work and my survival throughout.

First, and foremost, I need to thank (yes, thank!) my partner, Andrea Sch¨afer, for leading me along the path to the PhD. Somehow witnessing all the good and the bad moments during her PhD, ended up creating a positive image for me! Andrea also provided invaluable guidance and tips along the way, and created many shortcuts through the bureaucracy for me. I will never forget your assistance Andrea, and am deeply indebted to you. Vielen Dank, and may our love only grow stronger.

Another big Danke, goes out to our daughter, Moana Sch¨afer,who witnessed just over half of my efforts. Thanks for keeping my feet firmly planted on the ground and not letting me drift too far off into “PhD land”! Thank you for all our fun times, and my apologies for the times when my patience wasn’t sufficient to see your needs.

Naturally, I would like to thank the input from my supervisors of the years: to Stuart Wenham (UNSW) for his enthusiasm and encouragement; to Christiana Honsberg (UNSW and Georgia Institute of Technology, U.S.A.) for her moral and monetary support; to Francesca Ferrazza (Eurosolare S.p.A.) for the opportunities to see the “real” side of photovoltaics and for being a true friend; and to Jeff Cotter for his valuable advice during the latter stage of the thesis.

Several members brought their own special personalities to the Centre and made it a fun and challenging work place. These people include Keith McIntosh, Hamid Mehrvarz, Holger Neuhaus, Alex Slade, Bernhard Vogl, Rob Bardos, Matt Boreland, Martin Bruahart and Tom Puzzer. Thanks for all the ethical, moral and technical conversations. Thanks too for the great computer support, Laurie!

I would like to thank other people who assisted with TiO2 thin film characterisation: Dr. Tom Puzzer (UNSW) for SEM/AFM training; Prof. Robert Lamb (UNSW) and Dr. Matt Boreland (Toyota Technical Institute, Japan) for XPS analysis; Prof. David Jamieson (Univ. of Melbourne) for RBS analysis; Sally Rowlands and Prof. Trevor Redgrave (both Univ. of Western Australia) for training and access to the variable-angle spectroscopic ellipsometer (VASE); Dr. Alistair Sproul (UNSW), author of the forthcoming book “Ellipsometry for Dummies”; and to my father, Dr. Ray Richards (Lower Hutt, New Zealand), for his assistance in bringing me up to speed on thermochemistry analysis. 2

I am grateful for the guidance in my career provided by Dr. Andrea Sch¨afer,Prof. Mark Wain- wright, Prof. Stuart Wenham and Prof. Martin Green. The financial support provided by the Faculty of Engineering, the School of Electrical Engineering and the Centre for Photovoltaic Engineering was greatly appreciated. Publications Resulting from this Thesis (to date)

B.S. Richards (2004) Comparison of Dielectric Coatings for Buried-Contact Solar Cells: A Review, Progress in Photovoltaics 12 (in press).

B.S. Richards, S.R. Richards, M.B. Boreland, D.N. Jamieson (2004) High Temperature

Processing of TiO2 Thin Films for Application in Silicon Solar Cells, Journal of Vacuum Science and Technology A, 22(2): 339-348.

B.S. Richards, S.F. Rowlands, A. Ueranatasun, J.E. Cotter, C.B. Honsberg (2004) Reducing the Production Costs of Buried-Contact Solar Cells using Titanium Dioxide Thin Films, Solar Energy, 76(1-3): 269-276.

B.S. Richards (2003) Single-Material TiO2 Double-Layer Antireflection Coatings, Solar Energy Materials and Solar Cells, 79(3), 369-390.

B.S. Richards, S.F. Rowlands, C.B. Honsberg, J.E. Cotter (2003) TiO2 DLAR Coatings for Planar Silicon Solar Cells, Progress in Photovoltaics, 11(1), 27-32.

B.S. Richards, J.E. Cotter and C.B. Honsberg (2002) Enhancing the surface passivation of

TiO2 coated silicon wafers, Appl. Phys. Letters, 80(7), 1123-1125.

B.S. Richards, S.F. Rowlands, A. Ueranatasun, J.E. Cotter, and C.B. Honsberg (2001) Reducing the production costs of buried-contact solar cells using titanium dioxide thin films, Intl. Solar Energy Society Solar World Congress, 26-30 November, Adelaide.

B.S. Richards, J.E. Cotter, C.B. Honsberg and S.R. Wenham (2000) Novel Uses of TiO2 Films in Crystalline Silicon Solar Cells, 28th IEEE Photovoltaic Specialists Conference, Alaska, 375-378.

C.B. Honsberg, J.E. Cotter, K.R. McIntosh, S. Pritchard, B.S. Richards and S.R. Wenham, (1999), Design strategies for commercial solar cells using the buried contact technology, IEEE Trans. Electron Devices, 46(10), 1984-92.

B.S. Richards, J.E. Cotter, F. Ferrazza, C.B. Honsberg and S.R. Wenham (1998) Lowering the cost of commercial silicon solar cells, Proc. of the Environmental Engineering Research Event 1998, Avoca Beach, New South Wales, 303-308.

J.E. Cotter, B.S. Richards, F. Ferrazza, C.B. Honsberg, T.W. Leong, H.R. Mehrvarz, G.A. Naik and S.R. Wenham (1998) Design of a simplified emitter structure for buried contact solar cells, 2nd World Conference Photovoltaic Energy Conversion, Vienna, 1511-1514. Contents

1 Introduction 9

1.1MotivationforthisWork...... 9

1.2 Australia’s Solar Energy Resource ...... 11

1.3 Brief Theory of Solar Cell Operation ...... 12

1.4 Commercially Produced Silicon Solar Cells ...... 14

1.4.1 Screen-Printed Solar Cells ...... 14

1.4.2 Buried-Contact Solar Cells ...... 15

1.4.3 Buried-Contact Solar Cell Fabrication Sequence ...... 16

1.4.4 Simplified Buried-Contact Solar Cell ...... 16

1.5 Multicrystalline Silicon ...... 18

1.5.1 Issues with Multicrystalline Silicon ...... 19

1.6 Why use Titanium Dioxide? ...... 20

1.6.1 TiO2 Thin Films in Photovoltaics ...... 21

1.7 Thesis Overview and Goal ...... 25

2 Common Properties of TiO2 Thin Films 29

2.1Introduction...... 29

2.2PhysicalProperties...... 30

2.2.1 The Amorphous − Anatase − Rutile Phase Transformations...... 30

2.2.2 The Effect of Impurities on the Anatase − Rutile Phase Transformation 32

2.2.3 Substrate Type ...... 33

3 4 CONTENTS

2.2.4 Film Defects ...... 34

2.2.5 Film Density ...... 34

2.2.6 Non-Stoichiometric TiO2−x Thin Films ...... 36

2.3OpticalProperties...... 37

2.3.1 Refractive Index, Extinction Coefficient and Scattering...... 37

2.3.2 TiO2 vs. Single ...... 39

2.3.3 Variation with Deposition and Annealing Temperature ...... 40

2.3.4 Variation with Deposition and Annealing Ambient ...... 45

2.3.5 Variation with Other Deposition Conditions ...... 48

2.3.6 Optical Properties of Highly Porous TiO2 Films ...... 48

2.4ElectricalProperties...... 50

2.4.1 TiO2: Insulator or Conductor? ...... 50

2.4.2 Non-Stoichiometric TiO2−x Thin Films ...... 50

2.4.3 Variation with Deposition or Annealing Ambient ...... 51

2.4.4 Doped TiO2 ...... 51

2.5 Chemical Properties ...... 52

2.5.1 Chemicals used in Making Solar Cells ...... 52

2.5.2 HydrofluoricAcid...... 54

2.5.3 Other Acids and Bases ...... 56

2.6 Conclusions ...... 57

3TiO2 Thin Film Deposition Equipment 59

3.1Introduction...... 59

3.2 Overview of TiO2 Thin Film Deposition Methods ...... 62

3.3UltrasonicSprayDeposition...... 65

3.4 Theory of Ultrasonic Spray Deposition ...... 67

3.5 TPT: The TiO2 Precursor ...... 70

3.5.1 WhyTPT?...... 70 CONTENTS 5

3.5.2 The TPT→TiO2 Reaction...... 72

3.6DesignofUltrasonicSprayDepositionSystem...... 74

3.6.1 Selection of Ultrasonic Nozzle ...... 74

3.6.2 UltrasonicNozzlePerformance...... 76

3.6.3 Liquid Delivery ...... 77

3.6.4 Substrate Heater ...... 78

3.6.5 MotorizedStage...... 81

3.6.6 Spray Shaping ...... 82

3.6.7 Miscellaneous Equipment ...... 82

3.6.8 Operation of the TiO2 SpraySystem...... 85

3.7DesignofCVDSystem...... 87

3.7.1 Motivation...... 87

3.7.2 TPT Bubbler and Temperature Control ...... 87

3.7.3 Vapour Bubbler ...... 88

3.7.4 OperationoftheCVDSystem...... 88

3.8 Conclusions ...... 89

4 Characterisation of TiO2 Thin Films 91

4.1Introduction...... 91

4.2FTIRSpectroscopy...... 92

4.3RamanSpectroscopy...... 95

4.4X-rayPhotoelectronSpectroscopy...... 98

4.5RutherfordBack-ScatteringSpectroscopy...... 99

4.6 Ellipsometry ...... 101

4.6.1 Overview...... 101

4.6.2 Ellipsometers ...... 104

4.6.3 Lorentz Oscillator Model ...... 106

4.6.4 SurfaceRoughnessModel...... 107

4.6.5 Ellipsometric measurements of Spray Deposited TiO2 Thin Films . . 107 6 CONTENTS

4.6.6 SE measurements of CVD TiO2 Thin Films ...... 109

4.7 Reflectance Spectrophotometry ...... 109

4.8ElectronMicroscopy...... 109

4.8.1 USD-Deposited TiO2 Thin Films ...... 110

4.8.2 APCVD-Deposited TiO2 Thin Films ...... 110

4.8.3 CVD-Deposited TiO2 Thin Films ...... 110

4.9AtomicForceMicroscopy...... 115

4.10 Chemical Resistance ...... 117

4.11 Conclusions ...... 118

5 Enhancing the Passivation of TiO2-coated Wafers 121

5.1Introduction...... 121

5.2 Stability of TiO2 atHigh-Temperatures...... 123

5.2.1 Titanium Contamination of Silicon ...... 123

5.2.2 Reduction of TiO2 ...... 125

5.3 Methods of Achieving Surface Passivation with TiO2 Thin Films ...... 130

5.3.1 Growth of SiO2 at the TiO2:SiInterface...... 133

5.3.2 TiO2 onPSG...... 137

5.4 Conclusions ...... 139

6 Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source 141

6.1Introduction...... 141

6.2 TiO2 as a Phosphorus Diffusion Barrier ...... 142

6.2.1 Experiment...... 144

6.2.2 Results...... 144

6.2.3 Discussion...... 146

6.3 TiO2 asaPhosphorusDopantSource...... 149

6.3.1 Experiment...... 151

6.3.2 ResultsandDiscussion...... 151 CONTENTS 7

6.4 Conclusions ...... 152

7TiO2 Antireflection Coatings 153

7.1Introduction...... 153

7.2 Previous Developments in AR Coatings ...... 155

7.2.1 TheoryandDesignofARCoatings...... 155

7.2.2 TiO2 ARCoatings...... 160

7.2.3 AR Coatings ...... 165

7.3 Varying the Optical Properties of TiO2 ...... 167

7.3.1 DepositionTemperature...... 168

7.3.2 Annealing Temperature ...... 174

7.3.3 Deposition Ambient ...... 178

7.3.4 Annealing Ambient ...... 181

7.4 Development of Novel TiO2 ARCoatings...... 183

7.4.1 Single-layer TiO2 ARCoatings...... 183

7.4.2 Double-layer TiO2 ARCoatings...... 186

7.5 Performance of TiO2 DLAR-Coated Solar Cells ...... 193

7.6 Conclusions ...... 196

8 Conclusions 199

8.1 Summary ...... 199

8.2 Applicability to Various Solar Cell Processes ...... 203

8.3SuggestionsforFurtherWork...... 205

ATiO2 AR Coating Modelling Parameters 207

A.1 Variation of n and k withDepositionTemperature...... 208

A.2 Variation of n and k with Annealing Temperature ...... 212

A.3 Variation of n and k with Deposition Ambient ...... 218

A.4 Variation of n and k with Annealing Ambient ...... 220

A.5 TiO2 DLARCoatings...... 225 8 CONTENTS Chapter 1

Introduction

Photovoltaics (PV) will play an important role in the world’s future energy trends, however the major hurdles faced in widespread implementation of renewable energy are of a social, not technical, nature. This chapter briefly discusses Australia’s potential to become a major player in the future solar energy industry. The differences between the two dominant com- mercially produced silicon solar cells, screen-print (SP) and buried-contact (BC), are briefly examined. The BC technology was introduced in the mid-1980’s and designed for single crystal silicon (c-Si) wafers. A trend over recent years has seen lower-cost, multicrystalline silicon (mc-Si) wafers now dominate the marketplace. The BC fabrication sequence has several lengthy high-temperature processing steps fabrication. The processing costs currently outweigh the performance enhancement offered by the buried-contact technology when using mc-Si wafers and, today, only SP solar cells are produced on mc-Si substrates.

The most crucial high-temperature processing step in the buried-contact fabrication sequence is the growth of a thick, thermal layer. If this layer could be replaced by a thin di- electric film, deposited at low temperature and such a film was able to withstand processing in phosphorus-containing ambients, a buried-contact solar cell could be fabricated with about one hour of high-temperature processing steps. This proposition seemed attractive for ap- plication of the BC technology to multicrystalline silicon wafers in a cost-effective manner.

One logical choice for this film is titanium dioxide (TiO2), due to its prevalence in the PV industry, high optical performance and low cost. A review of roles that TiO2 films have played in the photovoltaics industry is presented, before discussing the overview of the remainder of the thesis.

1.1 Motivation for this Work

The amount of solar energy that strikes in a period of a few days is greater than the amount of fuel burnt overt the course of the whole human history,1 which encourages one to

9 10 1. Introduction think about ways in which this energy could be effectively harnessed in order to satisfy our ever-increasing demand for energy. While this may sound like a gross oversimplification of the impending energy crisis facing humankind, it brings into perspective the amount of power that is constantly being radiated by the sun. The primary focus of most developed nations needs at this stage should be looking at ways of radically reducing the current consumption of energy and raw materials. In fact, we are all so used to consuming that our language now defines our role in society as “consumers”. A 1995 study stated that for each American, 20 tonnes of new materials have to be provided every year, including energy equivalent to 7.6 tonnes of oil (or 12 tonnes of coal).2 If the world’s projected population in the year 2070 consumed energy at this rate, the world energy production would have to be fourteen times greater than its current capacity and all potentially recoverable energy resources would be depleted in about eighteen years2!

Therefore, if close to 11 billion people are to live on this in 70-years time we need to become “conservers” rather than consumers. The earth does not contain enough resources for each person in a developed nation to sustain our resource-intensive lifestyle. In the same manner, the economies of Western nations simply cannot continue to grow at targeted rate of 4% p.a., no matter what the politicians say. The Australian environmental thinker and activist, Ted Trainer, claims that the only type of economic growth that can be said to be truly ecologically sustainable is one that has a 0% growth rate per annum.2 Trainer also argues that “technological fixes” will not get us there either, as the main problems faced are social.2 So while the pursuit of renewable energy is a worthy cause, a truly sustainable solution can only be reached if we were to consume a small fraction of what we use in our affluent lifestyle today.

If photovoltaics (PV), the direct conversion of sunlight into electricity via solar cells, is to have a major and timely impact upon our current global predicament, the costs of production need to be significantly reduced before PV can compete head to head with fossil fuels. The cost of generating electricity directly from solar cells is slowly, but steadily, reducing. Figure 1.1 shows that the cost of purchasing PV modules is expected to reduce from today’s 1 US$4 /watt-peak (Wp)to

100.00 1978 1980

1985 1990 10.00 1995 1998 2000

1.00 1 10 100 1000 10000 Average selling price (1998 price US$/Watt) Average selling Accumulated shipments (megawatts)

Figure 1.1: Historical cost of purchasing PV modules from 1978 to 2000, and an extrapolation of the linear trend to increased shipments in the future (from Green1).

1.2 Australia’s Solar Energy Resource

Australia has recently implemented a PV rebate scheme has at both State and Federal level in an attempt to help Australia meet its target of generating an additional 9500 GWh (about 1%) of electricity from renewable energy sources by the year 2010. There is definite potential for a much greater expansion in PV in Australia, considering that many areas of Australia receive more than twice the amount of solar radiation than countries such as Germany (as shown in Figure 1.2), which already has a large installed PV capacity. It is also apparent from Figure 1.2(a) that the majority of Australians do not live in the sunny central and northern regions of the country. However, the cities of Perth, Adelaide, and all of the densely populated East Coast (Sydney and further north) still receive four or more sunshine hours a day of full sunshine (1 kW/m2). This is an excellent (and free) resource, which remains relatively untapped. Additionally, PV panels can easily be mounted on existing north-facing roofs and included into fa¸cades of office buildings. This means that power can be generated from within densely populated areas, obviating the requirement of setting aside large amounts of land for such purposes. Australia’s excellent research record in the field of solar energy along with its abundant resources make it uniquely positioned to become a major player in the future solar energy industry. One only needs to look to Denmark in the case of wind energy to see how a strong domestic market can to a small country dominating the world market. 12 1. Introduction

Figure 1.2: Maps indicating the level of solar radiation received on a horizontal plane in (a) Australia and (b) Europe (in units of kWh/m2).4

1.3 Brief Theory of Solar Cell Operation

A typical silicon solar cell is shown in Figure 1.3(a). The standard n+-p silicon solar cell has a shallow junction formed near the front surface, a front ohmic contact in the form of fingers and a busbar, and a full metal rear ohmic contact. Light that is absorbed by the solar cell generates an electron-hole (e-h) pair that is able to contribute to current flow from the device. For small photon energies, the majority of the current is generated in the base (about 300 µm thick), while photons of energy greater than 2.5 eV generate current from within the first 1 µm of silicon. This is the reason why the recombination velocity at the front surface can have such a profound effect on this high-energy photocurrent. For high-efficiency solar cells the effect of rear surface recombination also becomes an important design parameter.

The area-normalised current density J (mA/cm2), of the solar cell is given by V +JRs V + JR ( nkT /q ) s J = JL − J0 exp −1 + ,(1.1) Rsh 2 where J0 is the dark saturation current density (mA/cm ), JL is the light generated current density (mA/cm2), q is the electronic charge (1.602 × 10−19 C), V is the operating voltage, −23 k is Boltzmann’s constant (1.380 × 10 J/K), Rs is the series resistance, Rsh is the shunt resistance, n is the diode ideality factor, and T is the operating temperature (K). The special case, where the voltage is zero and J = JL is defined to be the short-circuit current density,

Jsc. Figure 1.3(a) shows the equivalent circuit of a solar cell, including series and shunt resistances.

The Jsc is limited by optical losses such as reflectance from the front surface, shading of the front surface due to the metal contacts, and transmission of lower energy light out the back of the cell. An additional source of current loss is due to minority carriers recombining at the surfaces before they are collected by the junction. 1.3 Brief Theory of Solar Cell Operation 13

J – Pmpp, Vmpp

Jsc Rs

J V V mpp, mpp

q(V-JRs )/nkT JL J0(exp -1) Rsh Current Density, Power Current Density,

+ Voltage Voc (a) (b)

Figure 1.3: (a) Equivalent circuit of the basic p-n junction solar cell, and (b) a typical I − V (solid) and P − V (dashed) curve of an illuminated solar cell.

The operating point where J = 0 mA/cm2 in Equation 1.1 defines the open-circuit voltage

Voc of the solar cell, as shown in Figure 1.3(b). The minimization of both bulk and surface recombination are important in order to maximise the voltage at maximum power point. Excellent surface passivation can be achieved by using a thermally grown

(SiO2) layer, as discussed in Section 5.3.1. Bulk recombination will always be higher in mc-Si wafers compared to high quality c-Si FZ wafers, due to the grain boundaries between the , although lifetime enhancement measures such as gettering and hydrogenation may improve the situation somewhat. In practice, the highest Voc’s of mc-Si and c-Si solar cells have been limited to 657 mV5 and 710 mV.6

The most efficient operating point of the solar cell is at the maximum-power-point voltage

Vmpp and maximum-power-point current Impp, as shown in Figure 1.3(b). At this point, solar energy is converted into electrical energy with an efficiency of I V η = mpp mpp , (1.2) Pin where Pin is the incident solar power. Another parameter commonly used is the fill factor (FF), which is a measure of the “squareness” of the I − V curve and is defined as I V FF = mpp mpp .(1.3) Isc Voc Recombination in the depletion region will lower a solar cells fill factor, as will a small shunt resistance or a large series resistance.

The interested reader is referred to the books by Green7, 8 and van Overstraeten and Mertens9 for a more detailed discussion on device physics and operation. 14 1. Introduction

1.4 Commercially Produced Silicon Solar Cells

1.4.1 Screen-Printed Solar Cells

The first reference to a commercial screen-printed (SP) solar cell production line can be found in 1975.10 Since this time the process has remained relatively unchanged. Figure 1.4(a) shows the structure of a screen-printed (SP) solar cell. As the name implies, the metal contacts in Figure 1.4(a) are formed by screen-printing a metallic paste through a mask. However, the contacts are the main limitation of the device, which has an efficiency of typically 12 − 13%. This is because of the paste’s poor conductivity, its poor contact resistance to silicon, the poor aspect ratios achieved, and the inability to reliably produce thin lines in production. The nett result is that the metal lines are much wider than desirable (150−200 µm). To still allow a significant fraction of the light (92%) to strike the front surface of the silicon, the fingers are spaced widely (about 3 mm) apart. The electrons are then required to travel a large lateral distance through the thin emitter region before reaching the metal contact. For this reason the emitter is heavily doped with n-type (negative) carriers, affording better lateral conductivity. This, however, gives the cell a very poor response to high-energy (-blue) light which is absorbed very close to the front surface. An (Al) or /aluminium (Ag/Al) paste is screen-printed onto the rear of the cell and fired in a furnace to form a p-type ohmic contact. The aluminium also creates a back- surface-field (BSF), which repels electrons that travel towards the rear of the cell instead of towards the junction. A titanium dioxide (TiO2) antireflection (AR) coating is deposited near the end of the process, increasing the amount of light absorbed by the silicon. The total amount of high-temperature (> 750◦C) processing involved in fabricating a SP solar cell is typically about 30 − 45 min. Modules fabricated today with SP cells typically cost about

US$3.50 /Wp.

Figure 1.4: Schematic diagrams of a (a) screen-printed solar cell and a (b) buried-contact solar cell (adapted from Green8). 1.4 Commercially Produced Silicon Solar Cells 15

1.4.2 Buried-Contact Solar Cells

The performance of commercial silicon solar cells was enhanced greatly through the devel- opment of the buried-contact (BC) solar cell at the University of New South Wales in the mid-1980’s.11, 12 The BC solar cell, shown in Figure 1.4(b), is currently commercially pro- duced in large volumes under license by BP Solar. Conversion efficiencies of greater than 20% have been achieved on laboratory scale cells at UNSW,11 and an independently mea- sured production BC solar cell had an efficiency of 16.7%.13 The fingers are only 20 µm wide but are 50 µm deep, the grooves being made either with a laser or a mechanical dicing saw. The cell is less shaded due to the reduced metal area (about 3%) on the front surface, and therefore the fingers can be placed closer together, permitting the use of a lightly-doped emitter and giving the cell an excellent response to blue light. Figure 1.5 provides a visual comparison of the metal contact areas of a SP and BC solar cell. It can be easily seen that the fingers of the BC solar cell are much finer and occupy a much smaller fraction of the solar cell front surface.

Figure 1.5: Scanned images of (a) a Solarex screen-printed solar cell on mc-Si, and (b) a BP Solar buried-contact solar cell on textured c-Si. The AR coatings used are titanium dioxide (TiO2) and silicon nitride for the Solarex and BP Solar cell, respectively (not to scale).

The drawback of the BC process is that although there are substantial materials costs savings relative to SP cells, the processing costs are higher, over 30% of which can be attributed to the high-temperature processing steps14 (see Section 1.4.3). In making the BC technology commercially-viable, BP Solar have removed the lengthy high-temperature oxidation step and have replaced it with an alternative dielectric, namely silicon nitride.15, 16 The silicon nitride is deposited using a low-pressure chemical vapour deposition (LPCVD) system.16 LPCVD silicon nitride is typically deposited at a temperature of about 700◦C,17 and is 16 1. Introduction capable of acting as a phosphorus diffusion barrier and metallisation mask as well as an AR coating.16 While the LPCVD silicon nitride film acts in many ways as a drop-in replacement for the thermally-grown silicon dioxide (SiO2) layer, the high deposition temperature means that the surface passivation benefits normally associated with silicon nitride are not realised in that process.16, 18

1.4.3 Buried-Contact Solar Cell Fabrication Sequence

Figure 1.6 describes the fabrication sequence for a single sided BC solar cell on a high quality float zone (FZ) c-Si wafer. Four high-temperature (> 750◦C) steps of varying length are involved in fabricating a BC solar cell:

i) a deep, high-quality emitter is formed by performing a light (low dose) n-type (phos- phorus, P) diffusion on the wafers. This creates the collecting junction in the p-type wafers. While this process is relatively short (15 min), the length of the following processes are in the order of hours each.

ii) a thick SiO2 layer is grown a) as a diffusion barrier to protect the lightly doped emitter from the heavy groove diffusion, b) to bond with atoms of the disrupted silicon lattice at the surface, thereby improving the surface passivation, c) to facilitate electroless metal plating of the front contacts, and d) to act as an AR coating, reducing reflection losses from the cell’s front surface.

iii) a heavy phosphorus diffusion (close to the solid solubility limit) in the grooves permits good electrical contact between the silicon and the metal.

iv) an evaporated Al film is sintered for a few hours to form an ohmic contact and create a BSF at the rear.

1.4.4 Simplified Buried-Contact Solar Cell

The simplified buried-contact (SBC) solar cell was developed in an attempt to reduce the number of high-temperature processing steps in the standard BC fabrication sequence.19, 20 In the SBC solar cell a single emitter and groove diffusion is performed, as opposed to the two separate diffusions in the standard BC process. This reduces the number of high- temperature steps by one. More importantly, by performing the homogeneous diffusion before the deposition of the AR coating, the limitations on the choice of front-surface dielec- tric layer are considerably relaxed. The dielectric film now only has to act as a metallisation mask and a good AR coating, and does not need to act as a phosphorus diffusion barrier. This opens the door for low-temperature deposited films such as plasma-enhanced chemical 1.4 Commercially Produced Silicon Solar Cells 17

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Figure 1.6: Fabrication sequence for a BC solar cell on a FZ silicon wafer. 18 1. Introduction vapour deposited (PECVD) hydrogenated silicon nitride (a-SiN:H) and APCVD-deposited

TiO2.

Cotter et al. performed PC1D21 modelling and determined that efficiencies greater than 16.5% could be achieved for a single emitter and groove diffusion with a sheet resistance greater than 40 Ω/2 as long as the front-surface recombination velocity (SRVf ) was kept 22 below 20000 cm/s. This condition is relatively easy to achieve with SiO2 passivation which 23 typically results in SRVf ≈ 1000 cm/s.

A necessary modification to the BC process has been the optimisation of the sintering step. The electrolessly-plated nickel (Ni) film is normally sintered at 350◦C in order to form a nickel silicide (Ni2Si) ohmic contact. Additionally, the (Ni2Si) layer may act as a diffusion barrier to the subsequently deposited layer at normal module operating temperatures. It was found that performing a 350◦C Ni sintering step with a 45 Ω/2 groove diffusion resulted in low fill factors (< 75%) and a open-circuit voltage (Voc) that is degraded by up to 40 mV.24 It has been postulated that the degradation is caused by small-area Schottky contacts between the metal and the p-type base.24 This is attributable to either the nickel punching-through the n-doped grooves during sintering or the lack of a diffused region in some areas of the grooves. By reducing the Ni sintering temperature to 250◦C this problem has been circumvented.24 This has resulted in the fabrication of 9 cm2 solar cells with a conversion efficiency of 16.9 − 17.1%, with a sheet resistances of 39 − 50 Ω/2 on 1 Ω cm -doped untextured CZ wafers.24 The use of a PECVD a-SiN:H film as an AR coating 2 enabled a respectable short-circuit current density (Jsc)ofupto34.8 mA/cm to be achieved. 25 These cells did not require any SiO2 passivation layer. The disadvantage of the current SBC process is that it is not nearly as robust as the standard BC fabrication sequence, and requires further refining before it is able to withstand the rigours of a production environment.

1.5 Multicrystalline Silicon

The cost of the silicon wafers alone represent the largest fraction in the cost breakdown of a solar cell. The cost of the growing the silicon ingot and cutting it into wafers contributes 46% towards the final PV module cost.26 This is not because silicon is a rare material, it is in fact the most abundant element on earth, however the purification and ingot growth processes are extremely energy intensive. Additionally, the PV industry is somewhat reliant on silicon scrap and off-cuts from the industry for their feedstock. The refining processes to obtain the necessary purity for good electrical performance and cutting the ingots into

350 µm thick wafers costs about US$1.50 /Wp. The price is not strongly linked to economies of scale, as the PV industry in the past has been able to rely on the much larger semiconductor industry for technological advances in crystal growth. Additionally. as the semiconductor industry has moved towards larger and larger wafer sizes - 300 mm is the current standard - 1.5 Multicrystalline Silicon 19 the PV industry has been able to purchase the outdated equipment relatively cheaply.

Only recently, with the strong growth of the PV industry, has dedicated equipment for PV technologies been developed in order to reduce the cost of silicon wafers. Firstly, there is an industry trend towards fabricating solar cells on thinner wafers. There are several handling issues in production to be overcome before high yields can be obtained on 150 µm-thick wafers. However, the French PV manufacturer, Photowatt, has demonstrated a high yield with 200 µm thick wafers on an automated production line.27 Secondly, different methods of casting silicon into multicrystalline silicon ingots have been developed.

1.5.1 Issues with Multicrystalline Silicon

Multicrystalline silicon (mc-Si) ingots are now used by many of the world’s leading solar cell manufacturers, including BP Solar (incorporating Solarex), Kyocera, Eurosolare, and Photowatt, to name a few. Mc-Si ingots of up to 65 × 65 cm and weighing 230 kg are now being grown.27 There are several advantages to such an ingot compared to the traditional 100−150 mm round Czochralski (CZ) wafers. Firstly, the cost of mc-Si wafers is on the order of 15% less than c-Si wafers grown by the CZ process. Secondly, the overall geometrical yield from a 150 mm diameter c-Si ingot is only about 66%.27 This is because the ingot must be trimmed into pseudo-square wafers to permit a higher packing density of the solar cells once they are laminated into modules. If this is not performed then module costs will increase due to the extra amounts of glass required to fabricate a module with the same power rating. In comparison, the wafer yield from a 65 × 65 cm mc-Si ingot is about 84%.

There are several disadvantages to using mc-Si wafers. Firstly, electrical properties of the material quality are slightly poorer. This is due to increased numbers of minority carriers recombining at the grain boundaries between the crystallites, before they can be collected by the junction. Secondly, due to the random orientation of the crystallites in mc-Si, the benefits achieved from the standard alkaline chemical texturing are minimal and a good AR coating is necessary to prevent large reflection losses. Thirdly, all mc-Si material is different, varying from manufacturer to manufacturer and from batch to batch. The material behaves differently under high processing temperatures (> 950◦C) and while the particular parameters used for lifetime enhancement processes, like hydrogenation and gettering, may work well for one material they may actually degrade the quality of another material. This indicates that a robust technology is required that can tolerate these kinds of variations in a production environment.

For screen-printed solar cells these disadvantages represent only a small loss in performance. The mc-Si wafers are still chemically textured, even though the benefits are minimal. Screen- printed solar cells fabricated on CZ c-Si substrates maintain a slight performance advantage over mc-Si wafers, by about 0.5 − 1.0 % absolute efficiency. The device efficiency is limited, 20 1. Introduction not by the substrate, but by the screen-print process itself, and therefore can be more cost- effective on mc-Si wafers. This fact, along with the reduced cost of mc-Si wafers has seen mc-Si screen-printed solar cells capture a major share of the PV market. Additionally, large amounts of money are being invested in new facilities that will fabricate tens of megawatts more of this product. Although investment in renewable energy is applauded, an opportunity is being passed up by many companies to invest in more efficient technologies that make better use of the silicon substrate. The upper efficiency limit for a mc-Si solar cell has been demonstrated at UNSW. A laboratory scale (1×1 cm) passivated-emitter rear-locally-diffused (PERL) structure fabricated on Eurosil P48 material (Eurosolare S.p.A., Italy) resulted in a conversion efficiency of 19.8%.28 This result indicates that a significant improvement margin exists for the 12− 13% efficient commercially-produced solar cells if the right technology can be found. This work investigated the applicability of the BC process to mc-Si wafers, using low-temperature deposited titanium dioxide thin films as the replacement dielectric coating.

1.6 Why use Titanium Dioxide?

There are several motivations for investigating titanium dioxide (TiO2) thin films in this work. TiO2 thin films are used currently as the PV industry standard AR coating on the vast majority of screen-printed solar cells. The important implications of this are, firstly, that the industry is familiar with the technology and will not be reluctant to adapt to fabrication processes based around TiO2 and, secondly, that the necessary deposition equipment is operating today on the factory floor. Thus, the development of a new silicon solar cell technology that included TiO2 processing steps could be readily adopted by the PV industry without the typical long lead-in time for a new technology.

TiO2 exists in nature as the rutile, anatase, and . Titanium dioxide of the rutile form is a relatively abundant ,29 however anatase and brookite are extremely 30 rare in nature. TiO2 thin films are generally amorphous for deposition temperatures ≤ 350◦C, above which anatase is formed. The most stable crystalline phase, rutile, is formed at temperatures greater than about 800◦C. The brookite phase is rarely observed in deposited thin films. The functional properties of TiO2 films, powders and ceramics are strongly dependent on the phase of the material. For many applications, the size of crystals that are present also alter the behaviour of the film. Typical properties of TiO2 include:

• Electrical: high electrical resistance - resistivities of 1014 Ωcm29 and dielectric constants of up to 180 are possible for rutile crystals.31

• Mechanical: high durability32 and high hardness.33

• Optical: very high refractive index - up to 2.70 − 2.71 (at a wavelength of 600nm) for rutile thin films33, 34 - and excellent transmittance in the visible region. 1.6 Why use Titanium Dioxide? 21

• Chemical: good chemical resistance and high chemical stability.35, 36

TiO2 powders and thin films are used in an extremely wide range of commercial applications and research areas, including:

• Powders: as a in , , , , and ; in washing powder, , , foodstuffs, pharmaceuticals, photographic plates, for creating synthetic gemstones; and as a catalyst.

• Thin films: for ultra-thin capacitors and MOSFETs due to its extremely high dielec- tric constant; as humidity and sensor due to the dependence of its electrical conductance on the gases present; as an and a material for waveguides due to its high refractive index; as a protective coating and resistant barrier; and as a photoanode in solar cells due its photoelectric activity.

1.6.1 TiO2 Thin Films in Photovoltaics

The use of TiO2 films has already been explored to a certain degree in the field of PV.

Antireflection Coating

Lord Rayleigh first observed the antireflection (AR) effect in 1887, and Bauer presented the first theoretical treatment based on interference effects in 1934.37 Since this time the theory and application of optical coatings have been well developed. The use of an AR coating for solar cells is a more recent application, beginning in the 1960’s with the advent of PV as a remote power source in space. For a good theoretical treatment of AR coatings, mainly on glass substrates, the reader is referred to Heavens,38 and some theory is also presented in Chapter 7.

Many references to TiO2 thin films being experimented with as a solar cell AR coating appeared in the early 1970’s.39–44 These early experiments were aimed at increasing the efficiency of space cells, which commonly employed (SiO) AR coatings. At a wavelength of 600 nm, the refractive index of silicon is 3.94, glass about 1.52, SiO about

1.9, and TiOx in the range 1.9−2.4 (it is difficult to know the exact stoichiometry of titanium dioxide thin films that are deposited via evaporation or sputtering, and therefore these films are denoted as TiOx). As discussed in Section 7.2.1, an AR coating with a refractive index of about 2.45 is optimal for achieving minimal reflection losses for a glass-encapsulated silicon solar cell. Thus, the use of TiOx films became widespread in the PV industry for providing better optical coupling of light into the silicon. Although TiOx films are more absorbing to short wavelength light, this is of little importance for SP solar cells, which already exhibit a poor blue response due to the phosphorus ”dead-layer” at the front surface.8 It should be 22 1. Introduction noted that the ethyl-vinyl-acetate (EVA) films used to encapsulate that silicon solar cells and the glass cover plates also exhibit significant absorption of short-wavelength light.45

Since the early 1970s, TiOx has been the main AR coating employed by the PV industry. Nearly all SP solar cell production lines use an APCVD- or spray-deposited TiO2 AR coat- ing.16, 19 Kern and Tracy provide an extensive review of early AR coatings for silicon solar 46 cells, focussing on TiO2. Several improvements have been made to the process, includ- ing optimisation for the firing-through of screen-printed contacts47, 48 and for deposition on 49 textured surfaces, and since 1994 TiO2 has been investigated for application in BC solar 16, 19, 20, 22, 24, 35, 50 cells. TiO2 thin films are also employed as AR coatings on glass, transmitting visible light while reflecting heat-producing IR radiation.51

Surface Passivation

Surface passivation is an extremely important design consideration for high-efficiency silicon solar cells, especially at the front surface where the majority of the light is collected. The predominant recombination losses in c-Si are via defect levels within the bandgap and the large number of non-saturated Si bonds at the surfaces dominate those defect levels. In order to reduce these recombination losses and achieve high conversion efficiency the surfaces must be electronically passivated, and, in the case of solar cells, the passivation scheme should be stable under ultraviolet (UV) illumination for at least 20 years.52 Two common and well- characterized methods for silicon surface passivation are thermal oxidation at temperatures ◦ of about 1000 CtogrowSiO2 and plasma-enhanced chemical vapour deposition (PECVD) of silicon nitride. Methods of achieving surface passivation with TiO2 thin films will be discussed in detail in Chapter 5.

Metallisation Mask

TiO2 (anatase) films have been used as a dielectric mask for preventing electroless nickel and copper plating from occurring on the front surface of solar cells. TiO2 thin films of about

70 nm in thickness have successfully replaced the 350 nm thick SiO2 employed in the BC solar cell fabrication sequence for use as a metallisation mask.20 The necessary film properties for this application include chemical resistance and a dense, continuous film with no pinholes. ◦ TiO2 thin films deposited by spray-deposition at temperatures above 400 C have satisfied these criteria.20 Research performed at the University of Konstanz (Germany) has shown that the use of PECVD deposited a-SiN:H thin films as a metallisation mask is problematic.53 As mentioned previously, BP Solar employ an LPCVD silicon nitride coating on their BC solar cell production line. Films deposited by LPCVD are typically very dense and are well suited to acting as a metallisation mask. The primary difference between silicon nitride films deposited by PECVD and LPCVD is that the former can contain as much as 30 at. % 1.6 Why use Titanium Dioxide? 23 . This, along with the significantly lower deposition temperature of PECVD films, results in the PECVD film having a significantly lower density.17

Ohmic Contacts

54 Thin films of TiO2 have been used to create ohmic contacts to p-type mc-Si solar cells.

The TiO2 layers were deposited in between layers of aluminium screen-print paste and p- type silicon. The screen-print paste was then fired (850◦Cfor5− 30 min), which allowed the aluminium to interdiffuse with the silicon, creating a titanium silicide (TixSiy) film in the process. With the above firing conditions, the TixSiy films yielded a low contact resistivity of 1 − 13 × 10−5 Ωcm2.

MIS Solar Cells

Metal-insulator-semiconductor (MIS) solar cells rely on quantum mechanical tunnelling through a very thin oxide layer, less than 2 nm thick, for carrier transport.7 This is made possible by the extreme work functions of the metal. The top contact can be a thin metallic layer (< 10 nm), which is essentially transparent to light. As this layer will have a high resistivity a thicker contact grid is required to transport the current. Alternatively, very fine (5 − 10 µm) and closely spaced (50 − 100 µm) metallic fingers can be used. In this case, carriers that are generated between the fingers can be collected by a nearby grating line before recombining.

The latter MIS grid structure was investigated using a 2 nm thick SiO2 layer, followed by a 55 100 nm thick TiOx layer. The TiOx was intended to act as an inversion layer, inducing a layer of minority carrier near the front surface. Although this apparently improved efficien- 56 cies, it was later found that the TiOx was acting as an accumulation layer on p-type silicon.

This resulted in further research being performed with pentoxide, although TiOx would be suitable for n-type solar cells.

Transparent Conducting Oxide Layers

In the PV industry, transparent conducting oxide (TCO) layers are most commonly used in thin-film solar cells where low current are present. The most common TCO is --oxide (ITO). However, more recently TCO layers have been employed for silicon wafer based heterojunction solar cells from Sanyo.57 These cells have a p-andn-type amor- phous silicon films deposited onto the front and rear of the n-type silicon substrate. A TCO layer is employed to reduce the front contact area required. There is always an electrical vs. optical trade-off with TCO layers, and even though a relatively high short circuit cur- 2 rent density is achieved (Jsc =36.7 mA/cm ) the spectral response curve exhibits significant 24 1. Introduction absorption at wavelengths less than 700 nm.57

The TCO layers typically have a sheet resistance of about 10 − 50 Ω/2. -doped

TiO2 films with this sheet resistance and a refractive index of 2.2 − 2.5 have been deposited as TCO layers.58 However, a film thickness of greater than 1.5 µm was required to achieve the 50 Ω/2 sheet resistance. With thinner films a logarithmic dependence of the resistivity on the film thickness was observed. Therefore, these films are too thick to act as an effective

AR coating. As discussed in Sections 2.4.2 and 2.4.4, oxygen deficient TiO2−x thin films can have resistivities down to 102 − 10−3 Ω/2, indicating their potential as a TCO.

TiO2/c-Si Heterojunction Solar Cells

Investigations have been performed on the addition of indium (In) to TiO2 to form p-TiO2/n- 59 60 Si heterojunction solar cells. The TiO2 films were deposited by a ”spray-CVD” process.

The In-doped films lead to an efficiency increase of about 40% over pure TiO2 films, resulting in an efficiency of 14.1% under 100 mW/cm2 (AM1) illumination. The open circuit voltage approached 650 mV and a high fill factor of 0.82 was achieved.

Gordon deposited doped TiO2 film as an electrode underneath a fluorine-doped tin oxide 61 (F:SnO2) layer to form a F:SnO2/TiO2/p-Si heterojunction solar cell. The patent deals mainly with niobium-doped TiO2, however other possible dopants that are also discussed include tantalum, , phosphorus, , and . The function of the TiO2 layer is to, firstly, overcome the “interfacial resistance” observed in SnO2/p-Si heterojunction cells, and, secondly, to act as an intermediate AR coating between the SnO2 61 62 and Si. The SnO2 has a refractive index of about 1.85 at 600 nm. It is noted that the

TiO2 does not exhibit sufficiently low resistance to act as a TCO. In both solar cells described here the TiO2 layer was about 100 nm thick.

The dye-sensitized solar cells described below are also heterojunction solar cells, however as these involve a liquid p-type electrolyte they will discussed in a separate section.

Dye-sensitized TiO2 Solar Cells

In 1991 a very different solar cell concept was presented based on dye-sensitized nanocrys- 63 talline TiO2 thin films and an iodine/iodide electrolyte. The sample structure is shown in Figure 1.7. The TiO2 is n-type while the dye is p-type. The cell works by conversion of photons to electrons by the dye and the subsequent transfer of electrons to the glass electrode by the TiO2 layer. The device had an efficiency of 7.1% under full sunlight, which increased to 12% under diffuse lighting. This solar cell has a large potential market due to drasti- cally reduced fabrication costs and conversion efficiencies that are comparable to amorphous silicon solar cells.64 The dye, in this case based, is used to photosensitize the 1.7 Thesis Overview and Goal 25

TiO2 film. A highly porous TiO2 layer with a large surface area to volume ratio is used to increase the amount of adsorbed dye. This increases the absorption properties of the device in the visible spectral region. Spray-deposition techniques have been used for depositing 65 the nanocrystalline TiO2 films and the highly porous, CVD-deposited films presented in Chapter 7 could potentially be used for this type of solar cell.

Figure 1.7: Sample structure of a photoelectrochemical cell using nanocrystalline TiO2. The arrows indicate the direction of light (from Li et al.64).

1.7 Thesis Overview and Goal

The aim of many silicon-wafer based PV research groups worldwide is to develop a new, commercially-viable, fabrication technology suitable for mc-Si wafers, in order to bridge the gap between the “five-year plan” for thin film dominance of the PV marketplace and the stock-standard product of the last 20 years, screen-printed solar cells.

The objective of this thesis project is to develop, understand and evaluate novel applica- tionsofTiO2 thin films to silicon solar cells. TiO2 is identified as an unique material with significant potential owing to its excellent optical and electrical properties. TiO2 thin films also appeared to be an attractive option due to the possibility of depositing them at a low- cost. While amorphous TiO2 thin films have a long history of being used as an AR coating on screen-printed solar cells, very few examples of the application of polycrystalline TiO2 (especially anatase) thin films to photovoltaics can be found in the literature.

Anatase TiO2 thin films exhibit a high refractive index and low absorption coefficient. This, along with their insulating properties and excellent chemical resistance, suggested 26 1. Introduction

that anatase thin films could be used as a direct replacement for the thermally-grown SiO2 layer in the standard BC technology. The replacement of the SiO2 layer with TiO2 promised much lower thermal budgets, simplified fabrication sequences and reduced processing costs.

The lack of literature regarding the behaviour and stability of TiO2 thin films under high- temperatures and different gas ambients meant that a significant amount of time was spent increasing this knowledge base.

In order for TiO2 to successfully replace SiO2 in the standard BC process, several key pa- rameters have to be explored. Firstly, it needs to be demonstrated that TiO2 thin films do not reduce the minority carrier lifetime of silicon wafers when processed at temperatures up ◦ to 1000 C. Secondly, TiO2 is known to be a poor option for passivating the surfaces of a silicon wafer, so therefore a successful method for enhancing the surface passivation of TiO2 coated silicon wafers needs to be developed in order to achieve high efficiencies. Fourthly, one crucial role of the SiO2 layer in the BC process is to act as a phosphorus diffusion barrier.

The performance of a TiO2 thin film in this role needs to be evaluated. Fifthly, an additional reduction of thermal budget is envisaged by combining the emitter diffusion and AR coating steps. In this manner, a TiO2 film doped with phosphorus atoms would be deposited onto the p-type wafer and, during a subsequent firing process, the phosphorus atoms diffuse out of the TiO2 and form an n-type emitter. Finally, as well as acting as a chemically resistant layer and an electroless metal plating mask, the optical performance of TiO2 AR coatings need to be optimised.

The primary goal is to develop a 16−17% efficient BC solar cell on planar mc-Si wafers. The application of the BC technology to textured c-Si wafers has been demonstrated by industry, however the current BC technology cannot be economically applied to mc-Si wafers due to the high-processing costs. Therefore, a simplified process, centred around using TiO2 as the thin dielectric film is sought. If TiO2 films prove to be successful then one expected outcome would be the evolution of a new solar cell technology that is readily applicable to today’s PV industry.

Following this introduction, Chapter 2 presents an extensive literature review, necessary to understand the physical, optical, electrical and chemical properties of TiO2. A review of thin film deposition techniques and a description of two deposition systems designed and constructed by the author are described in detail in Chapter 3. Due to the novel ways that TiO2 thin films were being implemented into solar cell processing sequences, extensive film characterisation was necessary to determine the variation of film properties with process conditions (Chapter 4). A novel method of overcoming the limited surface passivation achievable with TiO2 coated silicon surfaces is discussed in Chapter 5. Issues such as film stability at high temperatures and contamination are also addressed in this chapter. Novel PV applications of TiO2, such as its ability to act as a phosphorus diffusion barrier and phosphorus dopant source, are investigated in Chapter 6. A high-performance, commercially viable, double-layer antireflection (DLAR) coating, based on two or more TiO2 1.7 Thesis Overview and Goal 27

films with differing refractive indices, is demonstrated in Chapter 7. Modelling results of the performance of this DLAR coating on a planar BC solar cell will also be presented. Finally, Chapter 8 summarises the work, and presents opportunities for further research in the area. 28 1. Introduction Chapter 2

Common Properties of TiO2 Thin Films

The physical, optical, electrical and chemical properties of titanium dioxide (TiO2) depend greatly on the amorphous or crystalline phase of the material. TiO2 is a complex material with three crystalline phases, two of which are commonly observed in thin films - anatase and rutile. Anatase is commonly observed at film deposition temperatures of 350 − 700◦C, while higher temperatures promote the growth of rutile. Deposition temperatures lower than ◦ 300 C generally result in the formation of amorphous TiO2.

Amorphous TiO2 is a highest bandgap material (about 3.5 eV), and exhibits a low refractive index (about 1.9 − 2.0 for 600 nm wavelength light) and extinction coefficient. The chemical resistance of amorphous TiO2 films is poor in many acidic and basic solutions. Polycrys- talline anatase thin films, with an optical bandgap of about 3.2 eV, exhibit a much higher re- fractive index (as high as 2.532 at 600 nm for single crystal material) and a slightly increased absorption coefficient. With the crystalline structure comes increased chemical resistance, and dense anatase films are insoluble in many acids and bases. Rutile thin films (3.05 eV bandgap) have extremely high refractive indices (up to 2.70 at 600 nm for single crystal rutile) and below the bandgap absorption is still low. The chemical resistance of rutile is excellent, and after annealing at temperatures above 1000◦C it is insoluble in nearly all acids and bases.

2.1 Introduction

The aim of this work was to evaluate the performance of titanium dioxide (TiO2) as a drop-in replacement for the thick, thermally grown silicon dioxide (SiO2) layer in the buried-contact solar cell. The use of TiO2 immediately obviates one of the high-temperature processing steps, required to grow the SiO2 layer. TiO2 was chosen due to the ability of depositing films at low temperatures and at atmospheric pressure; the non-toxicity of the liquid precursor;

29 30 2. Common Properties of TiO2 Thin Films the familiarity of solar cell manufacturers with this film; and, their existing ownership of the necessary deposition equipment. It was anticipated that these factors would facilitate an easy transfer of a successful laboratory device into a commercial environment. However,

TiO2 is a relatively complex material, and three crystalline phases as well as the amorphous form of TiO2 exist. Since each of the these materials has different optical, electrical, and chemical properties it was necessary to perform an extensive literature review in order to predict how the TiO2 films would behave in different processing conditions.

This chapter will describe the properties of the crystalline phases most commonly observed in thin films, that of rutile and anatase, and amorphous TiO2. The third crystalline phase, brookite, is a less stable and common form of TiO2 is rarely observed in deposited thin films and will not be discussed here. There are many different parameters that affect the phase of a deposited TiO2 thin film. Some of these parameters are deposition method, deposition temperature, annealing temperature, deposition rate, deposition pressure, precursor type, reaction atmosphere, impurities present, and substrate type. The resulting phase or mixture of phases, plays a large role in determining the physical, optical, chemical, and electrical properties of the film.

This work relies greatly on the excellent optical properties of TiO2 thin films, as well as its chemical resistance and insulating properties. A summary of the physical, optical, electrical and chemical properties reported in the literature, with an emphasis on those relevant to solar cell fabrication, is presented in the following sections of this chapter.

2.2 Physical Properties

2.2.1 The Amorphous − Anatase − Rutile Phase Transformations

◦ 66, 67 Amorphous TiO2 thin films can be deposited at temperatures as low as 100 − 150 C.

Amorphous TiO2 does not have a strict crystallographic structure, often incorporates voids within the material, and has a relatively low density. For TiO2 thin films formed by chemical reaction, the lowest temperature crystalline phase of TiO2 that can be obtained is anatase.

To obtain polycrystalline anatase, the film can be either deposited as amorphous TiO2 and then crystallised by annealing at a higher temperature, or deposited as polycrystalline ma- terial directly. Nearly all published results indicate that the transition from an amorphous to anatase film occurs at about 300 − 365◦C, regardless of whether this is the deposition or annealing temperature. Rutile films are initially observed on silicon substrates at deposition temperatures above 700◦C, and more typically from 900 − 1100◦C. It should be noted that anatase is a metastable phase of TiO2, and the conversion to rutile involves a collapse of the anatase structure, which is irreversible.68, 69 Figure 2.1 indicates the structure of an anatase 2.2 Physical Properties 31 and rutile crystal. Although rutile and anatase are both of tetragonal crystallographic struc- ture, rutile is more densely packed and thus possesses a greater density.

Figure 2.1: Models showing the tetragonal structure of both anatase and rutile, and the denser structure of the rutile phase of the latter (adapted from Du Pont, Inc.70).

The TiO2 thin films deposited in this work are formed by chemical reaction, using chemical vapour deposition (CVD), and spray pyrolysis and systems. In this scenario, the substrate temperature is the primary means of controlling the deposited phase of the mate- rial. In contrast, physical vapour deposition (PVD) systems, such as evaporation, sputtering, and -beam deposition, the resulting phase and film structure is determined primarily by the kinetic energy of the impinging atoms. Therefore, the progression through the amor- phous, anatase, and rutile phases may not necessarily be expected. This is confirmed by the occurrence of rutile films at low deposition temperatures (< 450◦C) by carefully optimised deposition methods,34, 71 ion-assisted deposition,72 and reactive evaporation.73 However, the bulk of the discussion here pertains to TiO2 films formed by a chemical reaction, and where the substrate temperature dominates film growth characteristics.

Several researchers73–75 observed that the processing temperatures required to convert an anatase film into a rutile one are much higher than temperature required to deposit a rutile film directly. Agreement with this observation can be found in the literature, where the rutile phase is only present at low temperatures when it is deposited directly at that tem- perature.34, 71, 73, 74, 76–78 The formation of rutile at lower temperatures is facilitated by the kinetic energy possessed by the TiO2 during the deposition process, enabling the lowest energy state to be reached on the substrate. In contrast to the above observation, Fitzgibbons had previously claimed that the variation in physical and chemical properties of the films is determined solely by the maximum processing temperature, whether this be the deposition temperature or a subsequent annealing temperature.67 Amores et al. have published an excellent diagram indicating how the high-temperature sintering process and transformation of anatase to rutile crystals proceeds, shown here in Figure 2.2.79 The pro- 32 2. Common Properties of TiO2 Thin Films posed mechanism for the sintering and transformation of anatase into rutile involves several steps. Initially, the smallest particles (a) coalesce, forming bigger particles (b). The frac- tion of particles that are already large have been shown not to undergo sintering. The heat evolved from the exothermic sintering process causes the local nucleation of the rutile phase (c). Finally, as the conversion to rutile is also an exothermic process, this results in the transformation of the whole particle to rutile (d).

Figure 2.2: Proposed mechanism for the sintering and transformation of anatase into rutile. The smallest particles (a) coalesce, forming bigger particles (b). The fraction of particles that are already large have been shown not to undergo sintering. Heat evolved from the exothermic sin- tering process causes the local nucleation of rutile (c). The conversion to rutile is also an exothermic process, leading to the transformation of the whole particle to rutile (d) (adapted from Amores et al.79).

2.2.2 The Effect of Impurities on the Anatase − Rutile Phase Transformation

Many researchers have observed that the inclusion of a certain amount of impurities into

TiO2 can drastically alter the physical properties of the film. It has been shown that silicon and phosphorus inhibit the transformation from anatase to rutile, with 100% anatase phase being retained at temperatures as high as 870◦C for up to 3 hr for thin films80 and 1500 K for bulk samples.81 The retardation of the anatase-rutile transformation can be achieved with 3− 2− 3+ 80 68 68, 81 79 the following impurities: PO4 ,SO4 and Al ; AlPO4; SiO2; Co3O4 and MoO3; 82 + 83 84, 85 85 81, 86 Ce and Nb; K ; WO3; Na2O; and P2O5. Conversely, it is well known that other impurities enhance the formation of rutile at lower temperatures. These impurities 79, 85 79, 87 85 include CuO2; V2O5; and NiO, CoO, MnO2,Fe2O3.

Most researchers agree that oxygen vacancies are responsible for the overall transformation mechanism.68, 69, 88 Thus, the and fluorides (such as Li+1,Co+2 and Mn+4)thatassist 2.2 Physical Properties 33 the transformation can substitute for Ti+4 in the anatase lattice, resulting in the creation of −3 −2 oxygen vacancies. On the other hand, the inhibiting effect of other impurities (PO4 ,SO4 , Nb2O5) has been explained by the reduction of oxygen vacancies due to the substitution of Nb+5 and S+6 into the anatase lattice. Oxygen vacancies are also known to be created in hydrogen ambients, thereby favouring the transformation to rutile.69, 85, 88

It should be noted that during the growth of TiO2 thin films, contamination from the various chemical precursors can result. Titanium alkoxides are common TiO2 precursors, with the most frequently used being (also called tetraisopropyl titanate, TPT). The residue of the organic binders results in contamination of typically a few at.%, but as high as 13 at.%, being observed.64, 89–99 It is likely that carbon incorporation could be higher at low growth temperatures, as when higher temperatures were used the carbonate species decomposed, resulting in the removal of hydrocarbon fragments.64, 94, 99, 100 (TiCl4) is another common TiO2 precursor, and this results in con- tamination of the deposited film.91, 93, 101

2.2.3 Substrate Type

The effect of substrate type upon the deposited TiO2 film will be discussed only briefly, as nearly all experiments performed in this work employ silicon substrates. Several researchers have studied the properties of TiO2 films deposited onto various substrate types, including silicon, silica, quartz, alumina, titanium, copper, arsenide, stainless steel, as well as several types of glass. The different substrate types influence the physical properties of the deposited film, including the phase, texture or surface roughness. Optical and chemical properties also change, however these are primarily dependent on the phase and density of the polycrystalline TiO2 film. Possible reasons for the dependence of the TiO2 phase and properties on the substrate include the substrate’s surface conditions affecting the orientation and packing density of the molecules, and, with glass, the diffusion of metal into the film.102

Battiston et al. observed that anatase was the only phase present with films deposited onto titanium and stainless steel substrates and annealed at 750◦C.103 In the same work, an anatase-rutile mixed-phase was observed on fluosilicate glass substrates at 1100◦C and alumina substrates at 900◦C. At 1100◦C a single rutile phase was detected on the alu- mina. This is in agreement with other results, where either single-phase rutile or epitaxially- grown rutile films have been achieved upon alumina substrates.94, 104–107

Film depositions on glass substrates are typically limited to anatase, or anatase-rutile mixed phases due to the low (typically ≤ 650◦C) of most . However, results of depositions on quartz indicate a preference for the formation of anatase, even at temper- atures greater than 1000◦C.67, 108, 109 This is also true for glasses containing , where 34 2. Common Properties of TiO2 Thin Films it is possible for the Na+ ion to out-diffuse from the substrate and retard the formation of rutile.85, 102, 107, 110 In contrast, aluminosilicate glasses such as Corning 7059 are known 107, 110 to favour the formation of the rutile phase due to Al2O3 impurity. Out-diffusion of substrate elements has also been observed by Yuan and Tsujikawa, where up to 30% copper ◦ 111 concentration was found in a TiO2 film deposited onto copper sheet and fired at 800 C.

2.2.4 Film Defects

It important that the deposited films are relatively uniform in thickness, and do not exhibit pinholes. Several works have reported pinholes or similar defects in TiO2 films. Using scanning electron microscopy (SEM), Fitzgibbons observed an occasional pinhole at 20000× magnification in films deposited by CVD at 150◦C, however no pinholes were observed after annealing (300 − 1000◦C).67 Additionally, it was found that tensile stress caused the films to crack when the film thickness reached 400 − 500 nm. Nishide and Mizukami experimented with spin-coating of TiO2 films and found that some films exhibited 1 µm diameter pores when fired at temperatures below 500◦C, however these defects disappeared when the firing temperature was increased to 550◦C.112 This is in accordance with other results indicating 113 that TiO2 films deposited at lower temperatures are generally more heavily defected.

The existence of craters was noticed by Szlufcik et al. when screen-printing a titanium organo-metallic based , however the problem was alleviated with the addition on butanol 114 to the ink. Kern and Tracy commented that the existence of micro-pinholes and TiO2 particulates observed in pneumatically sprayed TiO2 antireflection (AR) coating did not im- pair solar cell performance.46 More recently, Golego et al. observed that some spray droplets can react on their way to the substrate, form a particulate and then become incorporated into the film.115 Spray depositions performed at very low temperatures (90◦C) resulted in the liquid layer cracking upon drying. Again, as the deposition temperature was increased slightly (to 120◦C) these defects disappeared. Kurtz and Gordon noted that by maintaining a large temperature differential between the substrate and the deposition equipment, avoids 116 TiO2 particulates from forming on the substrate.

2.2.5 Film Density

While both rutile and anatase possess a tetragonal crystallographic structure, rutile is more densely packed and thus possesses a greater density (4.26 g/cm3) than anatase (ρ = 3 70 3.84 g/cm ). For TiO2 thin films the highest observed density published to date is the 3 33, 117 range 4.09 − 4.10 g/cm for a rutile film. Amorphous TiO2 films exhibit a wide range of densities, from 2.4g/cm3 for porous films118 to more typical values of 3.2 − 3.65 g/cm3,119 while films deposited with a high kinetic energy have achieved densities in the range 3 33 3.6 − 3.8g/cm . It has been noted that the TiO2 films with a lower density can favour 2.2 Physical Properties 35 impurity diffusion.120

It is widely accepted that there is a linear relationship between density and refractive index 33, 67, 121–125 of a TiO2 thin film. The linear variation of the refractive index (measured at 550 nm) with density is shown in Figure 2.3 for films deposited by five different techniques.33 The equation of the line in Figure 2.3 is

nf =0.42751 ρ +0.91933 . (2.1)

This can be more usefully expressed for this work as n − 0.91933 ρ = f ,(2.2) 0.42751

3 where ρ and nf are the TiO2 film density (in g/cm ) and refractive index, respectively.

Figure 2.3: Experimental data from several researchers indicating that a linear correlation between TiO2 film density and refractive index is observed over for a wide range of values. Previously published data from Ottermann and Bange,121 Fitzgibbons et al.,67 Bendavid et al.,33 Hass34 and Ribarsky 126 was used. (adapted from Bendavid et al.33).

Additionally, the porosity of the film can be determined using the following relation127

n2 − 1 Porosity = 1 − f , (2.3) n2 − b 1 where nb is the refractive index of the bulk single crystal material. It should be emphasised that this value is an approximation due to the fact that, firstly, both anatase and rutile crystals exhibit strong birefringence and, secondly, that mixed anatase/rutile phases can exist. The values of mean refractive indices used in this thesis are 2.70 for rutile126 and 36 2. Common Properties of TiO2 Thin Films

34 2.532 for anatase, both measured at λ = 600 nm. Several researchers have noted that TiO2 films (anatase) derived from sol-gels tend to be highly porous in nature, sometimes up to 49%.128, 129

When TiO2 films are annealed at a temperature higher than their deposition temperature, particle sintering and crystallisation contribute to an increase in film density and refractive index, and, accordingly, a reduction in film thickness (see Wong et al. for example130). Generally, the type of film structure can be estimated before the deposition takes place, based on the Movchan-Demchishin structure zone model (SZM).72 Guenther expanded the SZM (see Figure 2.4) to include high density vitreous phases that can be achieved using PVD techniques.72 However, for films deposited via CVD or spray deposition, the structure of the film is predicted by dividing the substrate temperature (Tsub) during TiO2 deposition ◦ 131 by its melting point, which for TiO2 is Tmelt = 1832 C.

columnar dense vitreous porous dense polycryst. amorphous

0.3 0.4 1.0

Normalised substrate temperature (Tsub /Tmelt)

Figure 2.4: Structure zone model, expanded to include vitreous phases 72 observed with TiO2 (adapted from Guenther ).

2.2.6 Non-Stoichiometric TiO2−x Thin Films

Titanium dioxide thin films exhibit a bluish, purplish or greyish hue once they are reduced 29, 73, 88, 132–136 (become poor in oxygen) to TiO2−x or TiyOx. Oxygen deficiency can occur as a result of the deposition conditions. This is a common problem with evaporated thin

films, where the choice of source material (e.g., Ti, TiO, TiO2 or Ti3O5) and oxygen partial pressure of the system are critical. The formation of TiOx can also result from annealing TiO2 in a vacuum or hydrogen (reducing) ambient. In this case, oxygen is lost from the lattice to the furnace ambient. The optical and electrical properties of TiO2−x thin films will be presented in Sections 2.3.4 and 2.4.2, respectively. Non-stoichiometric thin films observed in this thesis will be discussed in Section 5.2.2. 2.3 Optical Properties 37

2.3 Optical Properties

2.3.1 Refractive Index, Extinction Coefficient and Scattering

The refractive index n of a material is primarily determined by the polarizability of the valence electrons.132, 137 Increased shielding of the positive nucleus results from elements with higher atomic weight, and this increases the polarizability of the electrons and consequently the refractive index. Silicon and are good examples with refractive indices in the infrared spectrum of 3.4 and 4.0. respectively. In compounds, the type of bonding also affects the refractive index, with highly covalently bonded compounds yielding higher refractive indices than predominantly ionically bonded compounds.137

The extinction coefficient k plays an attenuating role in the material. When the attenuation is solely due to absorption, it is termed the absorption coefficient 4πk α = . (2.4) λ Figure 2.5 displays a theoretical transmission spectrum for an optical thin film. It can be seen that a region of high transmittance (region II) is located between the region of short- wavelength fundamental absorption (region I) and the long wavelength limit (region III). The region of fundamental absorption is dependent on the electronic structure of the material, while absorption in the long-wavelength region is due to lattice vibrations and/or free carrier absorption. The transmission of region II is strongly linked to the stoichiometry and purity of the thin film.

The extinction coefficient of a film can also be increased by the scattering of light by surface and volume imperfections, such as surface roughness, porous microstructure, and density fluctuations due to crystallinity, and is thus dependent on the deposition method.132, 137 The term optical loss L is defined as being

L = A + S (2.5) with

1=R + T + L, (2.6) where A is absorptance, S is the scattered component, R is reflectance and T is transmit- tance. Wang and Chao noted that an increase in the extinction coefficient of amorphous 138 TiO2 thin films deposited by sputtering upon annealing. The drastic increase observed in k was attributed, firstly, to the formation of the anatase phase, and, secondly, to an increase in the surface roughness due to the polycrystalline nature of the annealed film. The scatter- ing loss of a rough surface is related to the root-mean-squared (RMS) surface roughness σ 38 2. Common Properties of TiO2 Thin Films

band-band transparent lattice vibration absorption region absorption

drop caused by free carrier impurity absorption absorption

I II III

Figure 2.5: Schematic of a theoretical transmission curve for an optical thin film. Three regions of absorption are shown, and each region has a different mechanism for optical absorption, as explained in the text (from Pulker 137). by139 4πσ2 TIS = , (2.7) λ2 where TIS is the total integrated scattering. Wang and Chao proposed that the extinction coefficient k could be divided into two components138

k = αa + αs , (2.8) where αa is the absorption coefficient (as defined in Equation 2.4) and αs is termed a scatter- ing coefficient. Figure 2.6 shows how the two contributions to the extinction coefficient, αa 138 and αs, vary as a function of annealing temperature in an oxygen ambient. It is observed that αa decreases with temperature, due to the reduction of oxygen vacancies in the film, while αs increases as surface roughness of the film becomes greater with the formation poly- crystalline anatase. If this evaluation was performed, say, at λ = 400 nm (near the bandgap) ◦ αa would increase at about 350 C due to lower bandgap of anatase (about 3.4 eV). It should be emphasised that Figure 2.6 was derived empirically for sputtered TiO2 films, and that wrinkling, cracking, or peeling of the film at temperatures above 300◦C may not occur for alternative deposition methods. For stoichiometric films deposited by spray deposition or

CVD, αa would not decrease at low annealing temperatures as there are no oxygen vacancies. ◦ At a temperature of about 800 C, αa would increase again due to the increase rutile fraction in the film (the bandgap of rutile is about 0.2 eV lower than that of anatase). The scattering coefficient αs would most likely level off with the increased rutile fraction, as the grain size 2.3 Optical Properties 39 for anatase thin films is typically an order of magnitude greater than that of rutile films (e.g., 20 nm versus 200 nm, respectively).119

Figure 2.6: Qualitative illustration indicating the behaviour of the ex- tinction coefficient k, and its absorption component αa and its scattering component αs with increased annealing temperature in an oxygen ambient (from Wang and Chao138).

2.3.2 TiO2 Thin Film vs. Single Crystal

The refractive index of a TiO2 thin film is typically much less than that of an anatase or rutile crystal, while the extinction coefficient of the deposited material will generally be greater than that of the bulk material. Figure 2.7 displays the refractive index for anatase single crystals. Due to the limited studies performed on anatase single crystals the data used for the anatase curves is shown in Table 2.1. Data for rutile is shown in Figure 2.7 for comparison.140 As both anatase and rutile are birefringent crystals it is necessary to calculate the mean refractive index nmean for a randomly oriented polycrystalline thin film. This is done using Equation 2.9 below

2n⊥ + n n = ,(2.9) mean 3 where n⊥ and n are for oscillations perpendicular and parallel to the optical axis, respec- tively. The mean refractive indices of 2.70 for rutile126 and 2.532 for anatase34 (both at 40 2. Common Properties of TiO2 Thin Films

λ = 600 nm) were selected as a reference bulk value to represent dense polycrystalline ma- terial.

Figure 2.7: Published values for the refractive index of single crystal anatase, taken from Meyer and Pietsch,141 Hass,34 Fitzgibbons et al.,67 Kingery et al.,142 Washburn,143 and Kim.144 The dispersive curve for single crystal rutile from Kim is also given.144

Very little anatase and rutile single crystal absorption data has been published in the visible as these materials are essentially transparent. The optical bandgap of anatase and rutile 145, 146 145 is about 3.2eV and 3.05 eV, respectively. The optical bandgap of amorphous TiO2 is commonly reported as being around 3.5 eV.147 Hence, optical absorption will increase with the successive transformations from amorphous to anatase to rutile material. It can be seen in Figure 2.8 that anatase has an absorption edge with a lower steepness, which is attributed to the presence of excitons and more imperfections and disorder in anatase crystals.120 Figure 2.9 shows the exponential dependence of the absorption coefficient at 10 K for different polarisations when illuminated with UV light.120

2.3.3 Variation with Deposition and Annealing Temperature

As previously discussed, many researchers have observed a direct linear relationship between 33, 67, 121–125 the refractive index and the density of the film. With TiO2 it is possible to deposit increasingly dense films, approaching the density of the bulk material, due to the mixture of amorphous, anatase, and rutile phases that can exist. Furthermore, while high refractive indices can be achieved for an amorphous film, a different set of deposition conditions can yield an anatase or rutile film with a higher refractive index.33 Thus, the optical properties 2.3 Optical Properties 41

Table 2.1: Refractive index data for anatase single crystal. The data for the anatase and rutile curves in Figure 2.7 are not presented in the table and can be found in Kim.144 The asterisked n values (∗) are for rutile.

Wavelength n n⊥ nmean Reference λ (nm) 435.83 2.7688 2.6576 2.732 Meyer and Pietsch141 546.07 2.5948 2.5161 2.569 589.30 2.5612 2.488 2.537 690.70 2.5097 2.4447 2.488

405 2.8760 2.7395 2.785 Washburn143 436 2.7688 2.6576 2.69467 492 2.6586 2.5691 2.59893 546 2.5955 2.5169 2.5431 578 2.5694 2.4950 2.5198 589 2.534 2.488 2.50333 623 2.5407 2.4709 2.49417 691 2.5106 2.4456 2.46727 706 2.5052 2.4409 2.46233

450 − − 2.703 Hass34 500 − − 2.615 550 − − 2.565 600 − − 2.532 700 − − 2.485

550 − − 2.57 Fitzgibbons et al.67

600 − − 2.52 Kingery et al.142

600 2.60∗ 2.90∗ 2.70∗ Ribarsky126

of a TiO2 film can be effectively tuned by adjusting the deposition temperature, bearing in mind that, in general, the extinction coefficient of a polycrystalline TiO2 thin film will be higher than that of an amorphous thin film.

Several works that have included extensive optical characterisation of TiO2 films will be re- viewed here. Hovel has published two excellent graphs demonstrating the trend of increasing refractive index with increasing deposition temperature, as shown below in Figures 2.10(a) 148 and (b). The TiO2 films were deposited by thermal spraying and the refractive index in Figure 2.10(a) was measured using ellipsometry at 633 nm.

Absorption values are reported in a variety of ways, including calculations of α or the TiO2 bandgap energy, and plots of optical transmittance or absorptance versus wavelength. As previously discussed, contributions towards optical losses in film arise from the fundamental 42 2. Common Properties of TiO2 Thin Films

Figure 2.8: Fundamental absorption edge of anatase and rutile single crystals, measured at a temperature of 10 K (adapted from Tang et al.120).

Figure 2.9: The exponential dependence of the absorption coefficient of single crystal anatase, measured at 10 K with light polarized in Ec and E⊥c directions (adapted from Tang et al.120).

absorption edge of the TiO2 film as well as surface and volume imperfections. Figure 2.11 indicates the general trend of increasing absorption coefficient with annealing temperature. This is in agreement with the formation of the anatase phase, which possesses a lower bandgap than amorphous TiO2. DeLoach et al. found that the absorption edge for sputtered

TiO2 films decreased from 3.41 eV for films with a small rutile fraction (< 0.17) to 3.22 eV for films with a large rutile fraction (> 0.7).149 2.3 Optical Properties 43

Figure 2.10: (a)Variation of TiO2 refractive index with annealing temper- ature, and (b) corresponding dispersive refractive index relations (adapted from Hovel 148).

Figure 2.11: Absorption coefficients of TiO2 films deposited at various temperatures (from Hovel 148).

Kamataki et al. performed optical characterisation of TiO2 thin films deposited by atmo- spheric pressure chemical vapour deposition (APCVD) at 300◦C in the wavelength range 250 − 850 nm.150 Measurements were also performed on samples that were annealed for 1 hr ◦ ◦ ◦ in both oxygen (O2) and (N2) ambients at temperatures of 500 C, 700 C and 900 C. Figure 2.12(a) shows a maximum in n at about 310 nm, while Figure 2.12(b) indicates that there is negligible absorption in the films at wavelengths greater than 350 nm. Both n and k exhibit a trend of increasing with higher annealing temperatures.

The optical constants of electron-beam (e-beam) evaporated TiO2 thin films were measured 44 2. Common Properties of TiO2 Thin Films

Figure 2.12: (a) Refractive index and (b) extinction coefficient of APCVD-deposited and annealed TiO2 thin films (adapted from Kamataki et al.150).

by Kim using spectroscopic ellipsometry (SE) in the spectral region 1.5 − 5.5 eV.140, 144 A double oscillator model for amorphous materials151 is used to fit n and k values. Kim also successfully modelled the film as being polycrystalline anatase with a 16% void content.140

Figure 2.13 displays n and k for the TiO2 evaporated film along with the “void-free” equiva- lent film. A three-layer model with varying void fraction was used to model the 96.4 nm-thick (total) film. The surface roughness was successfully modelled by 13.1 nm-thick top layer with a 34% void incorporation.

◦ ◦ Mardare and Hones deposited TiO2 thin films at 100 C (sample B) and 250 C (sample A) onto glass substrates using RF sputtering and determined n and k using SE.82 Figure 2.14(a) indicates that high refractive indices were achieved even at these low deposition temperatures, while Figure 2.14(b) demonstrates that the extinction coefficient remained below 0.02 for the sample deposited at 250◦C for wavelengths > 400 nm. The dispersive relations for Samples C and D in Figures 2.14(a) and (b) are for doped TiO2 thin films and will not be discussed here.

Szlufcik et al. investigated screen-printed TiO2 AR coatings and found that the refractive index increased linearly with annealing temperature up to a maximum of 2.30 at 800◦C.114 2.3 Optical Properties 45

Figure 2.13: Dispersive n and k values calculated from SE data for an e-beam evaporated TiO2 film (solid lines)and the equivalent “void-free” film (open circles). The refractive indices of polycrystalline anatase and rutile are also shown for comparison (from Kim144).

Figure 2.14: (a) Refractive index and (b) extinction coefficient of RF 82 sputtered TiO2 thin films. (from Mardare and Hones ).

2.3.4 Variation with Deposition and Annealing Ambient

Photovoltaic researchers have noted that the optical absorption of non-stoichiometric TiO2−x thin films increased in the short wavelength region, precisely where higher efficiency solar cells 46 2. Common Properties of TiO2 Thin Films

43 exhibited their improved response. As mentioned previously, non-stoichiometric TiO2−x films can be achieved under a variety of deposition and annealing conditions. The effect of 152 depositing TiO2 thin films in an oxygen poor environment is demonstrated in Figure 2.15. Films were evaporated at two different base pressures, 5 × 10−5 Torr and 1 × 10−4 Torr, and at one oxygen partial pressure. As shown in Figure 2.15, films evaporated without the presence of oxygen exhibited a much lower transmittance over the visible spectrum than the film deposited with an oxygen partial pressure of 5 × 10−5 Torr. The exact stoichiometry of the films is unknown, however a direct relationship between base pressure and transmittance can be observed.

Figure 2.15: Variation of transmittance of evaporated TiO2 films with base an oxygen partial pressure (from Jiao and Anderson152).

Zakrzewska et al. correlated the optical and physical properties of sputtered TiO2−x thin films deposited with different stoichiometries.153 The measure of stoichiometry was deter- mined as the fraction I/I0, which is the ratio of the titanium plasma line intensity during deposition to the 100% metallic titanium plasma line intensity. Higher I/I0 values indicate a greater departure from stoichiometry. The main graph in Figure 2.16 shows the increase in the film absorption coefficient with increasing I/I0. The inset graph in Figure 2.16 shows that the refractive index, measured at λ = 800 nm, increases linearly with decreasing of the TiO2−x film.

For TiO2 thin films deposited by ion-beam sputtering, it was found that the refractive index exhibited a maximum (2.52 at λ = 633 nm) with an oxygen concentration during 154 the deposition of [O2]=30%. The level of absorption in the film was found to decrease dramatically for [O2]< 30%, especially at the shorter wavelengths of 500 nm and 633 nm.

A difference in the optical properties of APCVD-deposited TiO2 films annealed in O2 and N2 ambients was noted by Kamataki et al., as shown in Figure 2.17. This was attributed to the growth of SiO2 at the TiO2:Si interface of O2 annealed samples. The analysis of Kamataki ◦ et al. determined that after 1 hr at 500 Can11.6 nm-thick SiO2 layerhadgrownatthe 2.3 Optical Properties 47

Figure 2.16: Variation of refractive index (inset) and absorption coeffi- cient (main graph) with oxidation state of the TiO2−x thin films (from Zakrzewska et al.153).

interface. This SiO2 growth rate is much greater than other researchers have observed at the same temperature.17

Figure 2.17: Effect of annealing ambient on the (a) refractive index and (b) extinction coefficient of APCVD-deposited TiO2 thin films annealed at 500◦C (adapted from Kamataki et al.150).

Fuyuki and Matsunami noted that the presence of a small amount of water vapour may 147 cause scattered values of the dielectric constant in CVD-deposited TiO2 films. Subse- 48 2. Common Properties of TiO2 Thin Films quent research demonstrated that the presence of water vapour altered the refractive index 155 from about 2.0 (no H2O) to about 2.15 (300 ppm H2O). Ardakani observed that as the hydrogen pressure in the laser ablation chamber increased, the measured reflectance of the 132 films decreased due to the higher concentration of oxygen-deficient phases of TiO2−x. ◦ Golego et al. determined that annealing spray-deposited TiO2 films in hydrogen at 500 C did not affect the absorption spectra, but an absorption peak at 400 − 600 nm appeared after the annealing temperature was increased to 900◦C.115 This indicates that absorption in the visible spectrum occurs after the film is reduced to TiO(2−x). Ottermann et al. deter- mined that there was a clear relationship between the hydrogen content of TiO2 thin films, deposited using different production processes, and their refractive index.156 An increase in refractive index was observed for decreasing hydrogen content, regardless of whether the films were amorphous or anatase. Films formed by chemical reaction methods, such as the sol-gel technique and dip-coating, exhibited high hydrogen contents. Thus, it is assumed that the organic molecules from the liquid precursor have not been fully decomposed by high-temperature processing.

Yoldas and O’Keeffe noted that spin-on TiO2 samples annealed in air had a refractive index of 2.1 (at λ = 546.1 nm), whereas the same coating annealed in vacuum had a refractive index of 2.4.157 This increase in refractive index was also observed for films annealed in or nitrogen, and was attributed to the increased densification resulting from the thermochemical reactions occurring with the residual terminal bonds in the absence of oxygen.102

2.3.5 Variation with Other Deposition Conditions

Depending on the type of deposition system used, there are many different parameters that may be tuned to adjust the optical properties of the TiO2 thin film. The results of Bendavid et al. are included here due to the extremely high refractive indices achieved - 2.56, 2.62 and 2.72 (at λ = 550 nm) for amorphous, anatase and rutile films, respectively. The films were deposited by filtered arc deposition (FAD) and a range of bias voltages were applied to the substrate. Figure 2.18 shows the refractive index and extinction coefficient of two TiO2 films deposited with different bias voltages, 0 V and −400 V, which were determined to be anatase and rutile, respectively.

2.3.6 Optical Properties of Highly Porous TiO2 Films

ATiO2 film with a high porosity, or low packing density, will have a reduced refractive index. The optical properties of porous films can be determined using a Bruggeman effective medium approximation (EMA), which can be expressed as158, 159

εm − ε 1 − ε (1 − fv) + fv =0, (2.10) εm +2ε 1+2ε 2.3 Optical Properties 49

Figure 2.18: Refractive indices and extinction coefficients of FAD- deposited TiO2 thin films. The film deposited with 0 V substrate bias is anatase and the −400 V film is rutile (from Bendavid et al.33).

where fv is the volume fraction of void in the film, and ε and εm are the dielectric functions of the unknown film and the main constituent material, respectively. The dielectric function of the void is taken to be unity, while the complex dielectric constant of a material is related to its refractive index and extinction coefficient by

ε =(n + ik)2 . (2.11)

TiO2 thin films deposited using the sol-gel method have exhibited void fractions in the 128 ranging from 25% at the exposed surface to 38% at the TiO2:substrate interface. This variation in void content translates directly into a variation in refractive index of the film, which, in this case, was 2.15 (at λ = 500 nm) at the outer surface and 2.05 at the inner surface.128, 160 The higher refractive index at the outer surface was attributed to densification and crystallisation processes that begin at the outer surface and gradually progress through the depth of the film during annealing.128 The void fraction observed in evaporated films 161 was slightly lower at 19.6 − 24.5%. Kim determined that TiO2 thin films deposited by e-beam evaporation had a void content of 16% when compared to dense, polycrystalline 140, 144 anatase. TiO2 thin films deposited by RF sputtering exhibited a much lower void content (< 5%),82 which can be attributed to much greater kinetic energy possessed by the atoms as the impinge on the surface.

Another issue related to porous films is the sorption (either adsorption or absorption) of water vapour. When the optical properties of a thin porous film are measured in a vacuum the results may be very different to those achieved in air. This is because on exposure to the atmosphere, the sorption of water vapour, which has a high refractive index than air n . 162 H2O =133, results in a marked increase in the refractive index of the film. Borgogno et al. measured the refractive index of a TiO2 thin film in vacuum and in air, and attributed the 0.1 higher n in air to the sorption of water vapour,163 while Leprince-Wang et al. and Nguyen Van et al. noted a 3% increase in n after exposure to air.161, 164 Ben Amor et al. 50 2. Common Properties of TiO2 Thin Films observed that it was possible to detect the adsorbed water vapour using Fourier transform infrared (FTIR) .165

2.4 Electrical Properties

2.4.1 TiO2: Insulator or Conductor?

TiO2 is an n-type semiconductor with a relatively wide bandgap of 3.05 − 3.5 eV. Single 13 crystal titanium dioxide TiO2 has a resistivity of about 10 Ω cm at room temperature, and about 107 Ωcm at 250◦C.29, 116, 155 These values are similar to conductivities reported for single crystal rutile:141 at 30◦C the conductivity was 5 × 10−14 Ω−1 cm−1 while at 260◦C −9 −1 −1 it had decreased to 3.3 × 10 Ω cm . Therefore, TiO2 is generally considered to be an insulator at temperatures less than 200◦C.141, 166

There are a wide number of applications for highly insulating TiO2 films, including its use as a super-thin gate dielectric in MOSFET devices.167 However, the electrical properties of the TiO2 film can be altered to become highly conductive for various applications, such as humidity and gas sensors,83, 168 or as an corrosion-resistant electrode.132, 169 The properties of TiO2 thin films have been likened to those of oxide (ZnO), with typically a large bandgap, small grain size, low carrier density, n-type conductivity, long photoconductivity relaxation, and low mobility.115 This section will briefly discuss the relationship between the electrical properties of TiO2 and the deposition parameters, annealing ambients, and dopant atoms.

2.4.2 Non-Stoichiometric TiO2−x Thin Films

The semiconducting properties of non-stoichiometric titanium dioxide are very dependent on the extent of the oxygen deficiency in the film.132, 170 The conductivity mechanism in undoped

TiO2 relies on a deficiency of oxygen atoms in the material. These oxygen deficiencies behave like n-type defects, with a typical density of 1018−1019 cm−3.171 As the stoichiometry departs from the ideal Ti:O ratio of 1:2, several changes in the film result. Firstly, as previously mentioned in Section 2.2.6, the TiO2−x films exhibit a blue, purple or black colour depending on the film stoichiometry. Secondly, due to the increased number of oxygen vacancies, the optical absorption in the visible is dramatically increased for these films (see Section 2.3.4).

Additionally, these sub-oxides, especially TiO and Ti2O3, are known to exhibit metallic conduction or semiconducting properties.58, 172, 173 The conductivity is observed to increase −10 −1 −1 dramatically with slight departures from stoichiometric TiO2, with < 10 Ω cm for −1 −1 −1 −1 −1 2 −1 −1 TiO2.00;10 Ω cm for TiO1.9995;0.8Ω cm for TiO1.995;and10 Ω cm for 141 −3 −1 −1 TiO1.75. Rao et al. determined that the resistivity of Ti2O3 is about 94 × 10 Ω cm 2.4 Electrical Properties 51

174 at room temperature. Feuersanger noted that the deposition of TiO2 thin films by direct evaporation of a TiO2 source results in the loss of O2, and the films were semiconducting 66 rather than insulating. Tsuda et al. showed that Ti2O3 exhibits a sharp transition in conductivity at 227◦C, exhibiting metallic conductivity above that temperature and semi- conducting below it.173

Thus, TiO2−x thin films would seem to be incompatible with high-efficiency solar cells, as the increased optical absorption is undesirable in an AR coating the buried-contact fabrication sequence requires an insulating film to act as a metallisation mask. All films used in this work were intentionally stoichiometric.

2.4.3 Variation with Deposition or Annealing Ambient

Much work has been performed determining the electrical properties of as-deposited TiO2 films, and the effect of the deposition or annealing ambient. It is commonly reported that thin films deposited in a low oxygen partial pressure result in the formation of sub-oxides such 113, 135 as TiO and Ti2O3. This results in increased electrical conductivity in the deposited films.175

When stoichiometric TiO2 films are annealed in a vacuum, hydrogen, or low oxygen pressure ambient, loss of oxygen from the film results.116,132,169,176 However, at room temperature it has been demonstrated that the film conductance does not increase when placed in a vacuum chamber and dry hydrogen gas is introduced.177 Yeung and Lam found that very low conductivity (10−13 Ω−1 cm−1) films could be achieved by performing an extended oxidation 178 of up to 30 hours. Stoichiometric, as-deposited TiO2 thin films possess a negative fixed charge density, in the order of −5 × 10−8 Ccm−2.117,178,179 Erkov et al. observed that after annealing in N2O this fixed charge density became zero, while annealing in vacuum created a 117 film with a positive fixed charge density. In comparison, SiO2 has a positive fixed charge density.179, 180

Takahashi et al. demonstrated that the resistivity of TiO2 thin films could vary by more than an order of magnitude depending on the measurement ambient.170 The lowest resistivity was observed in a hydrogen (reducing) ambient, consistent with the properties of an n-type semiconductor. Additionally, the film conductivity was observed to increase by 104 when illuminated with a fluorescent lamp.

2.4.4 Doped TiO2

58 Doped TiO2 films have been used as transparent conducting oxides (TCO), for creating heterojunction solar cells,59, 61 for suppressing leakage currents in capacitors,166, 181 increasing the sensitivity of gas detectors115, 168 and as electrodes for photoelectrolysis devices.63–65 52 2. Common Properties of TiO2 Thin Films

Dopant atoms have also been introduced into TiO2 thin films to create a diffusion source for an underlying silicon layer. These atoms then diffuse into the silicon during a subsequent high-temperature anneal. This will be discussed in more detail in Chapter 6.

The majority of dopants enhance the n-type semiconducting properties of TiO2. These dopants include niobium (Nb),116, 168, 182 tantalum (Ta),116 vanadium (V),116 fluorine (F),116 and hydrogen (H).116, 172 Dopants that change the film to being p-type include aluminium (Al),116 (Fe),116, 182, 183 and indium (In).59 There are instability issues with both hydrogen-doped (n-type) and all p-type TiO2 thin films. In the former case, the hydro- gen is very mobile and is likely to result in electrically unstable devices.116 In the latter case, the electrical behaviour of the Al- or Fe-doped film depends on the oxygen concentration, and it is possible for the film to revert to n-type.116, 182 Additionally, dopants such as alu- minium, iron,170 and ,59 are known to extend the photoactive response from the TiO2 film under visible light. Table 2.2 includes some references to published works on both doped and undoped TiO2 crystals and films, and their corresponding resistivity.

2.5 Chemical Properties

2.5.1 Chemicals used in Making Solar Cells

In general, the resistance of thin films to chemical attack is a highly desirable property for most applications.137 Schr¨oder stated the requirements for a film to be considered chemically resistant. A chemically resistant film must36

i) be insoluble to the attacking medium

ii) be impervious to the attacking medium, and

iii) form a solid bond to the substrate, preventing chemical attack at edges or through defects in the coating.

A large variety of chemicals are used to fabricate solar cells, but this is dependent on the specific type of solar cell and whether it is produced in a laboratory or industrial environment. A summary of the typical chemicals used at different stages in fabricating buried-contact solar cells include:

• (NaOH), hydroxide (KOH), or a mixture of hydrofluoric

(HF) and nitric (HNO3) acids: for texturing (basic solutions only) or etching silicon

• hydroxide (NH4OH), hydrochloric acid (HCl) or sulphuric acid (H2SO4): for removing organic and metallic contaminants from the surfaces of the wafers − each 2.5 Chemical Properties 53

Table 2.2: Measured resistivity of anatase and rutile crystals and thin films after annealing in various ambients and/or being doped with various impurities (adapted from Kurtz and Gordon116).

Resist. Sample Preparation Reference (Ω cm) Single Crystals - Undoped 1014 Stoichiometric rutile Clark29 ◦ 184 1 Rutile reduced at 1000 CwithanO2 partial Gautron et al. pressure of 10−11 Pa. ◦ 185 3 Rutile reduced at 700 Cfor2hrinH2, result- Butler ing in a carrier conc. of 4 × 1018 cm−3 2 −3 ◦ 10 − 10 Rutile reduced in H2 or vacuum at 750 Cfor Ginley and upto2hr Knotek186 −1 1 187 10 / 10 Anatase: as-grown/after repeated O2 anneal- Forro et al. ing Single Crystals - Doped 6 188 3 × 10 TiO2 doped with 1 mol % boron Johnson 4 188 2 × 10 TiO2 doped with 1 mol % phosphorus Johnson 3 188 3 × 10 TiO2 doped with 1 mol % niobium (Nb) Johnson 189 1.5 TiO2−xFx (x=0.002) Subbarao et al. 20 −3 190 0.1 TiO2 doped with 10 Nb atoms cm Bogomolov et al. 184 0.3 TiO2 doped with 1% Nb and reduced at Gautron et al. ◦ 1000 CwithanO2 partial pressure of 10−11 Pa Thin Films - Undoped 10 191 4 × 10 Evaporated films annealed in O2 for up to Yokota et al. 14 hr 1010 − 109 MOCVD at 200◦C as-deposited Fuyuki and Mat- sunami147 107 / Reactively sputtered anatase as-deposited / Tang 192 5 × 10−2 reduced at 450◦C in vacuum 6 4 −3 −6 135 10 − 10 Evaporation of Ti under 10 − 10 Torr O2 Chen et al. 107 − 102 1.6 − 17.2 µm thick films (higher deposition Takahashi et al.71 temperature to higher resistivity film) 103 Films deposited via spray-CVD method at Badawy and El- 400◦C Taher193 Thin Films - Doped 7 3 170 10 − 10 CVD of Al-doped TiO2 films, deposited at Takahashi et al. 435 − 480◦C 4 3 170 10 − 10 CVD of Fe-doped TiO2 films, deposited at Takahashi et al. 415 − 470◦C 2 170 4 × 10 CVD of Cr-doped TiO2 films, deposited at Takahashi et al. 435◦C 54 2. Common Properties of TiO2 Thin Films

of these chemicals is mixed with hydrogen (H2O2) and de-ionised (DI) water

at a ratio of 1:1:5. With NH4OH:H2O2:H2O this is referred to as an ”RCA1” clean, 194 while HCl:H2O2:H2O is called an ”RCA2” clean.

• HF: for removing silicon dioxide (SiO2) and doped SiO2 layers − diluted to about 5% in DI water

• Nickelex − commercial product from Transene Inc. (MA, U.S.A.) that contains nickel chloride. It is used for forming a thin electrolessly plated nickel layer in the grooved contacts in buried-contact solar cells.

• Enplate Cu-704 − another commercial product (Enthone, Melbourne) that has three

components (A, B, M), containing copper sulphate (CuSO4), (HCHO), NaOH, potassium cyanate (KCN), as well as other organic wetting agents. It forms the main bulk of the conductor in the grooves.

• Silver plating − contains (AgCN) silver cyanate. It is used as a flash coating to improve solderability

During the silicon etching procedures, the TiO2 film would not normally be present on the wafer, however the TiO2 would ideally be resistant to all of the chemicals used in the subsequent cleaning and metal plating processes. As the majority of results for the chemical resistance of TiO2 films are against HF this acid will be dealt with in a separate section.

2.5.2 Hydrofluoric Acid

A critical step in the buried-contact solar cell fabrication sequence is the removal of the

P2O5:SiO2 from the grooves before performing electroless metal plating, using either dilute hydrofluoric acid (HF) or buffered HF (1:15, HF:NH4F). This allows nickel to plate to the heavily doped silicon, however if even a thin oxide is present the metal will not plate. In general, the crystalline phases of TiO2, anatase and rutile, are much more chemically resistant than amorphous TiO2. Kurtz and Gordon noted that TiO2 films deposited via atmospheric pressure chemical vapour deposition at 400 − 600◦C were very chemically inert and that surpassed the chemical resistance of glass to attack by common solvents and acids.116 These films could not be removed mechanically or chemically, and etching away the glass substrate in HF left an intact TiO2 film.

Feuersanger reported that films deposited at 150◦C were easily etched in 10% HF.66 Fitzgib- bons observed that films deposited by CVD at 150◦C etched at 50 − 75 A/s˚ in 0.5% HF and very rapidly in 48% HF.67 Upon subsequent annealing at 350◦C, dilute HF undercut the

TiO2 layer by dissolving a thin interfacial SiO2 layer, while concentrated HF etched the film slowly and unevenly. Annealing at 700 − 1000◦C made the films resistant to dilute HF, and 2.5 Chemical Properties 55 concentrated HF resulted in undercutting. Spray deposited films annealed for 30 s at 450◦C were removed in 2 − 5 min when subjected to 1 − 5% aqueous solutions of HF,46 however the aim of this work was to ensure the TiO2 films were amorphous. Yokozawa found that TiO2 ◦ (anatase) films deposited by CVD at temperatures of 530 − 700 CinN2 hardly dissolved (< 0.03 A/s),eveninconcentratedHF.˚ 118 On the other hand, films deposited at the same temperatures, but in an N2 and O2 ambient, were less crystalline and etched at 4− 60 A/s˚ in diluted HF. Fuyuki et al. deposited TiO2 film deposited by metal-organic CVD (MOCVD) with water vapour present. The etch rates of these films in 10% HF were about 5 nm/min at 400◦C and 200 − 1000 nm/min at 200◦C.155 Brown and Grannemann used buffered HF to ◦ remove TiO2 films annealed in O2 at 1000 C, by undercutting grown SiO2 and leaving the 31 ◦ TiO2 film intact. Rausch and Burte observed that films deposited at 450 C by low-pressure CVD (LPCVD) etched in 50% buffered HF.195 Frenck et al. determined that parameters other than in situ and ex situ temperature have only a minor influence on etch rate in buffered HF.124 If no post-deposition annealing step is performed, the PECVD films are only chemically resistant to buffered HF with deposition temperatures greater than 270◦C. Lee noted that the film etch rates in buffered HF decreased from 2.9 nm/s to 0.14 nm/s as the anatase fraction of the film increased from 0.35 up to 0.7.113

Thus, there is a trend of increased chemical resistance of TiO2 thin films deposited or an- nealed at higher deposition temperatures. The majority of films treated at temperatures greater than 300◦C either could not be etched, or only etched slowly, in HF. This tempera- ture coincides with the amorphous−anatase transformation.

There are a number of published works that contrast the above trend of increased chemical resistance with increased annealing or deposition temperatures. Rausch and Burte claim ◦ that the etch rate of TiO2 films annealed at 900 C was four times greater than the as- deposited layer, which etched at a rate of 10 nm/min in 50% buffered HF.195 However later in the paper, it is stated that the annealed TiO2 film lifted off, exhibiting therefore a greater chemical resistance. Secondly, Harbison and Taylor reported an etch rate of 700 A/min˚ in 48% HF for material grown by hydrolysis at 800 − 1000◦C.179 This anomalous result does not agree with the above trend, especially considering the high growth temperature. Other ◦ results of TiO2 films deposited at relatively high temperatures (550 − 650 C) that etched, include the films prepared by Balog et al. using MOCVD, which etched in a solution of 5% HF.196 also found that films formed by rapid thermal annealing (30 s at 550 − 650◦C) of titanium metal at these temperatures could be etched in HF.197 It is claimed that the films are polycrystalline rutile, which is in agreement with the high dielectric constants reported, but not the poor chemical resistance. It has been noted that the chemical resistance to HF 141 is highly dependent on the film deposition technique and water-content of the TiO2 films.

The inclusion of a small fraction of SiO2 (6.25 at. %) into the TiO2 film results in rapid etching (70 nm/min) in buffered HF, even after annealing at 1260◦C for 10 min.198 56 2. Common Properties of TiO2 Thin Films

2.5.3 Other Acids and Bases

Kurtz and Gordon found that TiO2 films were chemically resistant to attack by common solvents and acids, and surpassed the chemical resistance of glass.116 The chemical resistance 141 of TiO2 to sulphuric acid (H2SO4) is very dependent on the film preparation technique, but water-free TiO2 is insoluble in all other acids and bases. This is in agreement with

Barksdale, who observed that TiO2 is known to be slightly soluble in H2SO4, HF, and a few strong alkalis, however after annealing at 1000◦C it is almost completely chemically inert.199

◦ TiO2 films prepared by Balog using MOCVD at 550 − 650 C were able to be etched in a 196 solution of 70% H2SO4. Fitzgibbons found that the chemical resistance increased with temperature, with films deposited at 150◦C etching at 25 − 40 A/s.˚ 67 However, even films ◦ annealed at 1000 C still etched slowly (1000 A/hr˚ in boiling pure H2SO4. At the same an- nealing temperature, TiO2 films were observed to etch very slowly in 85% H3PO4. Yoldas and O’Keeffe found that for 1 wt. % concentrations of H2SO4,H3PO4, and HNO3 no observ- able deterioration was observed after 75 days.157 That research also showed an increase in ◦ chemical resistance to acids for TiO2 films annealed at higher temperatures (up to 400 C). ◦ TiO2 films deposited by the sol-gel technique and baked at 120 C were observed to etch in 200 boiling 0.3 mol HNO3 (pH=0.5).

◦ Schr¨oderfound that an observable colour change with TiO2 films baked at 200 C occurred after about 4 hours immersion in a 10% HCl solution, whereas after baking at 550◦Cthe same colour change occurred after more than 200 hours.36 Szlufcik et al. determined that a ◦ 1 wt. % HCl solution completely deteriorated TiO2 films fired at 300 C after 26 days, however no observable deterioration was achieved after firing at temperatures greater than 600◦C.114 For spray-deposited films (450◦C) no observable deterioration was found after placing the films in a 1% solution of boiling HCl for 1 hour.46

The resistance of TiO2 thin films to various bases has also been reported in the literature.

Schr¨oder performed experiments with TiO2 thin films in a 10% NaOH solution, and found that for films annealed at 200◦C an observable colour change occurred after 5 hr.36 Increas- ing the annealing temperature to 500◦C delayed the same colour change observation to more ◦ than 200 hours. TiO2 films baked at 120 C were badly corroded after placing them in boil- ing 0.5 mol NaOH solution (pH=13.5) for 30 min.200 The addition of 1 at. % boron oxide

(B2O3) significantly increased the chemical resistance of the film. Kern and Tracy found ◦ no observable deterioration for spray-deposited TiO2 films annealed at 450 Cina1%so- 46 lution of NH4OH. For screen-printed TiO2 films the chemical resistance to 1 wt. % NaOH ◦ 114 and NH4OH solutions was poor for annealing temperatures of 600 C. No observable de- terioration was detected after increasing the annealing temperature to 800◦C. Honsberg et al. found that cleaning 60 nm-thick, spray-deposited TiO2 films on silicon wafers in RCA solutions resulted in thickness reduction in the film of 6 nm.20 This is most likely due to the 194 NH4OH in the RCA1 clean. Yoldas and O’Keeffe found that the chemical resistance of 2.6 Conclusions 57

◦ TiO2 to 1 wt. % NH4OH films fired in a vacuum at 500 C) was greater than that of films fired in air.157 For the air-baked films, the deterioration was obvious after 7 days, and complete after 10 − 20 days. NH4OH based etches can also be used for removing unreacted titanium in films.113

Changes in the reflectance spectra of TiO2 films containing small fractions of SiO2 that were fired at 400◦C were noted for 1% HCl solutions after 24 hours.201 No change was noted for a

1% NH4OH solution after 192 hours, but the 1% NaOH solution resulted in the swelling of the film and a reduction in refractive index. The hierarchy of chemical attack for these films 36 was given as NaOH > HCl > NH4OH. This is in agreement with the results of Schr¨oder, but 157 in contrast with the NH4OH results from other research. However, Yoldas and O’Keeffe postulated that the increased chemical resistance to sol-gel films baked at only 80◦ could be due to the retention of organic groups.157 The results of Schr¨oder indicate that the chemical 36 resistance of TiO2 containing SiO2 is significantly poorer than pure SiO2 films.

◦ TiO2 films deposited at above 400 C have also been shown to be resistant to the nickel and copper electroless metal plating solutions used in the buried-contact solar cell fabrication sequence.20 The chemical resistance to the copper plating solution is significant as it is strongly basic (pH=11) and contains NaOH.

Thus, increased deposition and annealing temperature result in greater chemical resistance for the majority of other acids and bases - a similar trend to that observed with hydrofluoric acid.

2.6 Conclusions

The properties of TiO2 depend greatly on whether the sample is bulk material or a thin film and the phase of the material. With CVD techniques the phase is usually directly related to ◦ the substrate temperature, with amorphous TiO2 forming at temperatures less than 300 C, the metastable crystalline phase of anatase in the temperature range 300 − 700◦C, and the stable crystalline phase of rutile at temperatures greater than 700◦C. Various impurities and substrate types are known to partially or totally impair or enhance the transformation of aTiO2 thin film to rutile. A useful linear relationship can be found between the refractive index and TiO2 film density.

While optical constant data for rutile is fairly easy to find, it was necessary to collect small sets of data for anatase from over many decades in order to construct a dispersive refractive index model. Anatase has a fundamental absorption edge with a lower steepness than rutile, due to the increased disorder observed in anatase crystals. A trend of increasing refractive and extinction coefficient with increasing deposition temperature is commonly observed.

Thus, accurate control of the temperature results in accurate tuning of the TiO2 thin film’s 58 2. Common Properties of TiO2 Thin Films refractive index and, to a certain extent, extinction coefficient.

The electrical properties are briefly discussed. Although TiO2 is a wide bandgap n-type semiconductor, at room temperature it behaves like an insulator. The film resistivity is extremely sensitive to the deposition method and on the availability of oxygen to the system. Films deposited in oxygen poor ambients exhibit greatly increased electrical conductivity and optical absorption, however subsequent furnace oxidation processes can reduce these oxygen vacancies. Doping of TiO2 thin films in order to increase the electrical conductivity or photoconductivity of the film has been experimented with in many different applications.

The chemical resistance of TiO2 thin films change markedly during their amorphous- polycrystalline transition. Amorphous TiO2 film are highly soluble in hydrofluoric acid, while dense, polycrystalline films can be insoluble. TiO2 films seem to be most susceptible to etching in strong basic solutions such as sodium hydroxide and ammonium hydroxide. The chemical resistance to sulphuric acid is dependent on the film preparation technique. There is a definite trend of increased chemical resistance to all chemicals with increased film deposition or annealing temperature. Chapter 3

TiO2 Thin Film Deposition Equipment

After evaluating the desired film properties and performing a literature survey on the possi- ble deposition methods, the author designed and constructed two TiO2 thin film deposition systems. The first system used an ultrasonic atomisation spray nozzle in order to create an aerosol of the TiO2 precursor. The reasons for choosing ultrasonic spray deposition (USD) and the TiO2 precursor, tetraisopropyl titanate (TPT) are discussed. A diagram of the sys- tem is presented and the necessary components described. With this system, very dense ◦ TiO2 films could be deposited at a temperatures of 450 C and the very shallow deposition angles successfully prevented TiO2 film deposition in grooves scribed in the front surface of the wafer. Thickness uniformity and chemical resistance problems with the electroless metal plating solutions, used for contact formation in the buried-contact solar cell fabrication se- quence, arose due to the frequent inclusion of TiO2 particulates (1 − 30 µm in diameter) in the 70 nm thick films.

For this reason the TPT was placed in a stainless steel bubbler, resulting in the development of a simple atmospheric pressure chemical vapour deposition (CVD) system. TiO2 films deposited using the CVD system exhibited much greater thickness uniformity and a lack of particulates. Additionally, it was also possible to deposit films anywhere in the temperature range 150 − 450◦C, enabling the refractive index to be tuned. The tradeoff with the new system was that the film density decreased significantly.

3.1 Introduction

The previous chapter described the physical, optical, electrical and chemical properties of

TiO2 thin films, primarily as a function of deposition or annealing temperature. Based on previous experience with a TiO2 deposition system at UNSW, TiO2 thin films with the

59 60 3. TiO2 Thin Film Deposition Equipment following properties were desirable for this work:

• Physical: Dense, defect-free films with a thickness of about 70 nm. The thickness uniformity should be within ±10%. Films that are dense and lack defects (such as

large, amorphous TiO2 particulates) will be impervious to chemical attack. Denser films will typically perform better as a diffusion barrier as well.120

• Optical: Dense TiO2 films will also enable a refractive index of about 2.4 for light of 600 nm wavelength to be achieved at low deposition temperatures (< 450◦C). This value is the optimum refractive index for an antireflection (AR) coating for a silicon solar cell encapsulated under glass. An excellent AR coating is necessary to reduce the amount of reflected light from the planar multicrystalline silicon (mc-Si) wafers.

• Electrical: The TiO2 film must be insulating in order to act as a metallisation mask to the electroless metal plating solutions. This requires that the films are stoichiomet- ric, and do not exhibit the oxygen vacancies that form the conduction mechanism in

reduced TiO2−x thin films.

• Chemical: The TiO2 films should be polycrystalline anatase or rutile in order to with- stand the necessary wet chemical processing during solar cell fabrication. The films need to withstand RCA cleaning, dilute hydrofluoric acid, and the electroless metal plating solutions.

A simple TiO2 pressurized spray system had already been in operation for some years at the UNSW, using tetraisopropyl titanate (TPT) as the precursor.202 This system used a spray-painting gun pressurized with nitrogen, and the TPT flow was controlled via a needle valve. A TPT aerosol was formed and transported to the wafer by the nitrogen flow. The wafer was held by vacuum onto a stainless steel block that was placed on top of a 2 kW stove element. The whole deposition process was carried out inside a fumecupboard. This system was used for several years, and it demonstrated that geometry could be used to direct the emerging aerosol, so that when spraying at shallow angles the TiO2 could be kept out of the grooves.202, 203

There were a number of problems with the TiO2 films and the deposition system. Firstly, it took over an hour to deposit a 70 nm thick TiO2 film onto a 2”-diameter wafer. This was not due to the low TPT flow rates, although the needle valve was being operated at its limits to avoid a too high a flow of TPT from emerging from the nozzle, but because of the cooling effect of the aerosol on the stainless steel block. After several seconds spraying, the temperature of the block dropped by over 30◦C, and it was necessary to wait a few minutes before continuing spraying. This highlighted the second problem, which was that the heater was very inefficient at transferring heat to the stainless steel block, and this excess heat made it uncomfortable for the user. Thirdly, there was no way of controlling the relative humidity, 3.1 Introduction 61 and the relative humidity on rainy or overcast days (up to 65%) made spraying impossible due to the formation of white TiO2 particulates in the film.

A replacement TiO2 system was to be designed and constructed by the author. In order for the TiO2 thin films to exhibit the desired properties, listed above, the requirements for the new system were determined as follows:

• Temperature: Significant fluctuations in the temperature at about 400◦C could result

in the TiO2 film having a mixed amorphous-anatase phase. Naturally, this outcome is undesirable, so a target of achieving a maximum variation of substrate temperature of ±10◦Cwasset.

• Relative Humidity: Excess humidity will result in large TiO2 particulates sticking to deposited film.202, 204 Therefore, a necessary feature of the system was to have adjustable and repeatable humidity control. It was anticipated that mixtures of dry

nitrogen (N2) and wet nitrogen (N2+H2O) could be fed into the system in order to control the relative humidity. A meter should display the current relative humidity in the system.

• Deposition Time: To achieve accurate control of the film thickness, a deposition time of about 5 − 10 min per wafer was decided upon. The film thickness would be judged visibly.

• Deposition Area: As this project had commercial relevance, it was desirable that a 4”-

square wafer could be TiO2-coated, although the research would typically be performed on 2”-round and 5 × 5 cm-square wafers.

• Geometry Control: To prevent TiO2 from being deposited in the laser-scribed grooves,

film deposition was to occur on the very top surface only. Any TiO2 that entered the grooves would inhibit electroless metal plating. The pressurized spray-system at UNSW had successfully fulfilled this goal by spraying a TPT aerosol onto the wafers at an angle of 10 − 20◦.

• Substrate Heater: A dedicated substrate heater is required to efficiently transfer heat into the substrate. The temperature should be accurately controlled by dedicated electronics.

• Automation: The previous pressurized spray system required the user to manually operate a spray-gun. In the new system it was deemed necessary that the user should only need to load and unload the wafer from the system, the user’s hands remaining free during film deposition.

• Safety: The system should be enclosed so that it can safely operate on a standard laboratory bench. It also needs to have exhaust facilities. It was desirable to continue 62 3. TiO2 Thin Film Deposition Equipment

using TPT as the precursor as it is a safe, non-toxic liquid, and obviates the need for expensive gas handling systems.

• Cost: There was a budget of A$10,000 for the new TiO2 deposition system.

Initially, a new system was designed using an ultrasonic atomisation nozzle. Figure 3.1 shows the USD system in its final form, including the ultrasonic nozzle and generator, syringe pump, cartridge heaters embedded in the stainless steel block, motorized sample stage, and temperature controller. Some components have been omitted for clarity, including a nitrogen heater, air-knife heater, motor speed controller, both the oxygen and humidity sensors, as well as regulators and flow meters to accurately control gas flows.

The following section will describe the literature review performed to evaluate simple, flexible and cost-effective techniques for depositing TiO2 films. Subsequently, the theory of ultrasonic spraying will be discussed, and a thorough description of the necessary system components will be presented.

3.2 Overview of TiO2 Thin Film Deposition Methods

Titanium dioxide has been deposited by many different techniques, including

• hydrolysis and pyrolysis,49, 60, 65–67, 71, 74, 78, 115, 148, 168, 177, 179, 193, 205–207

• pneumatic spraying,46, 208

• ultrasonic spraying,65, 106, 206, 209–212

• dip coating100, 109, 111, 156, 157, 200, 207, 213

• plasma enhanced chemical vapour deposition (PECVD),93, 97, 113, 121, 124, 134, 214–216

• atmospheric pressure chemical vapour deposition (APCVD),116, 130, 150, 204, 217–221

• metal organic chemical vapour deposition (MOCVD),58, 71, 77, 92, 94, 96, 103, 105, 106, 108, 147, 155, 167, 170, 195, 196, 211, 220–227

• ultra-high vacuum chemical vapour deposition (UHV-CVD),228

• low pressure chemical vapour deposition (LPCVD),90, 91

• evaporation,31, 34, 73, 76, 119, 121, 137, 156, 161, 213, 229–231

• spin-on methods,89, 100, 112, 121, 157, 198, 200, 232, 233

• sputtering,119, 120, 145, 149, 154, 165, 234–236 3.2OverviewofTiO2 Thin Film Deposition Methods 63 r PID contolle temperature C C q q r 445 450 DC 1.8 W 12 V ultrasonic generato line switch (x2) switch thermo -couple vacuum st. steel block relay housing TPT storage TPT bottle rings vacuum motor 3-way valve TPT heaters cartridge nozzle 2 ultrasonic hot N syringe air-knife syringe pump syringe 0.02 ml/min

Figure 3.1: TiO2 ultrasonic spray deposition system. 64 3. TiO2 Thin Film Deposition Equipment

• ion assisted deposition,72, 107, 140, 237, 238

• plasma anodisation,113

• reactive ion plating,121,156,213,231

• laser ablation,132, 239

• filtered arc deposition,33

• atomic layer epitaxy,240 and

• screen-printing.114, 241

A primary consideration is that the growth morphology, crystalline structure and stoichiom- 73, 110, 122, 237, 242 etry of TiO2 thin films are very sensitive to the deposition conditions. This a disadvantage for many physical vapour deposition methods, such as evaporation, where a large variation in the observed optical properties arise from only a small changes in the de- 107, 161, 164 position conditions. Therefore, the need for stoichiometric TiO2 films with minimal absorption suggested that a deposition method where the film stoichiometry is controlled by a chemical reaction would enable more consistent results.

The chemical reaction of a TiO2-precursor to form TiO2 can either proceed by hydrolysis or pyrolysis. With hydrolysis systems, separate gas lines of nitrogen or argon are bubbled through heated baths of a liquid TiO2 precursor and water. The two delivery lines are then brought together close to the substrate where the reaction takes place. Pyrolysis systems are similar, except that a water bath is not required as the TiO2 precursor decomposes upon reaching the heated substrate. Both of these systems have the advantage of simplicity, although they may be relatively inflexible, as tubing diameters have to be designed around predicted flow rates.

To keep the deposition system as simple as possible, maximise throughput, and keep costs at a minimum, systems with a vacuum chamber were not considered to be a viable option. This excluded evaporation, sputtering, and the majority of CVD systems. However some exper- iments were performed with an APCVD system,204 owned by Eurosolare S.p.A. (Nettuno, Italy). The APCVD system is designed around a belt furnace, and, as the name implies, the depositions are performed at atmospheric pressure, so no vacuum is required. The system 217 is capable of depositing TiO2 onto 580 solar cells per hour, which corresponds to about a 5 MW yearly throughput for screen-printed solar cells.

Spin-on methods were not seriously considered, as the throughput of any such system would be limited in a production environment. Screen-printing is commonly used in the PV in- dustry for depositing metallic contacts, about 30 − 50 µm thick, to solar cells. Szlufcik et 114 al. demonstrated that screen-printing could also be used for depositing TiO2 thin films.

As the thickness of the metallic contacts are about 500 times thicker than the TiO2 thin 3.3 Ultrasonic Spray Deposition 65

films it is not known how the thinner films behaved with regard to reproducibility, squeegee wear, and thickness uniformity. Following the screen-printing of the organometallic ink, the samples were fired in a three-step process of 30 min duration. The firing is a relatively slow process as first the thick film needs to settle for 15 min to obtain a uniform film, with subse- quent drying performed at 125◦C for 5 min. The final crystallisation was performed in a belt furnace at temperatures between 500◦C and 900◦C. Lengthy drying procedures are required to remove the substantial amount of organic solvents added to the TiO2 precursor. This was also necessary in the pneumatic spraying technique used by Kern and Tracy.46 Kern and Tracy developed a system for production was developed that was capable of coating 4500 cells per hour (around 30 MW per year), based on batch processing. Each batch took 30 s to receive a coating, followed by three separate heating steps to remove organic groups in the film. It has been noted by other researchers that the temperature required to crys- tallise an amorphous film is significantly greater than the temperature needed to grow such a crystalline film.74 Therefore, to lower thermal the budget and processing costs it would be desirable to deposit a polycrystalline TiO2 thin film in one step without subsequent heat treatment steps.

3.3 Ultrasonic Spray Deposition

35, 65, 106, 206, 210, 212 Several researchers have used USD for depositing TiO2 and other thin films.209, 211, 243 Blandenet et al. describes a method for depositing many different metal- lic oxides based on the pyrolysis of an aerosol.210 If an ultrasonic beam is focussed on the surface of a liquid the vibrations result in the formation of on aerosol. The chemical to be sprayed is contained in a glass container attached to a high frequency generator that vibrates at 800 − 1000 kHz. The chemical precursor for TiO2 films was butyl orthotitanate diluted in acetyl- and butanol. Air or nitrogen is passed through the glass container, transporting the aerosol close to the heated substrate, which is subsequently decomposed by pyrolysis. Other researchers have also used commercially available ultrasonic nebulizers,65, 211 while more recently ultrasonic atomizing nozzles have been used.35, 58, 106, 212, 243 Versteeg et al. used an ultrasonic nozzle to inject small quantities of a TiO2 precursor into a vacuum chamber, maintained at a pressure of 0.1−1 Torr.106 It is stated that the low pressure cham- ber facilitates film uniformity over large areas. Liang implemented an ultrasonic nozzle as a modified injector in an APCVD system.58 DeSisto and Henry deposited oxide thin films by USD.243

The primary advantages of USD are that, firstly, there is a very narrow size distribution of the droplets in the aerosol. Secondly, by altering this droplet size, the droplet→solid reaction mechanism can also be changed. This is also dependent on gas flows and the nozzle to substrate distance. For example, smaller droplets will have already decomposed by pyrolysis and will strike the wafer as a solid, while larger droplets will not have had time to vaporise 66 3. TiO2 Thin Film Deposition Equipment completely. Alternately, the spraying conditions can be adjusted so that the majority of the aerosol is a gas when it contacts the wafer, resulting in a process similar to CVD.209, 210, 244

Additional advantages of USD that were potentially relevant for this work, included:

i) The potential flexibility of the system and excellent parameter control. This included low flow rates, variation of spray angle, variation of power to atomize liquids of different viscosities and at different flow rates, and control over the shape of the spray plume using air-jets.245, 246

ii) The possibility of spraying at a shallow angle to keep the TiO2 out of the grooves, scribed on the front side of a buried-contact solar cell. In this manner, the simple geometry of the system prevents the spray from entering the grooves.20, 202, 203 With other deposition systems a lot more care would need be taken to keep the grooves

free of TiO2. This is especially true for the majority CVD systems which provide a conformal coating of the surface.

iii) Since the distribution of the droplet sizes can be varied by choosing the excitation frequency (see Section 3.4), it was postulated that larger droplet sizes (greater than the groove width) would not enter the grooves, but would still coat the front surface of the wafer.247

iv) It is possible to pulse-feed the liquid into the nozzle, enabling very low flow rates to be achieved.106, 245

v) It is possible to shape the spray plume by directing jets of gas across it. This can be used for focussing the spray plume or to spread it out.245

vi) As decomposition by hydrolysis or pyrolysis is a chemical process, this would en-

sure that stoichiometric TiO2 was deposited. The deposition of TiO2 by physical vapour deposition methods, evaporation and sputtering for example, can result non-

stoichiometric TiOx films. TiOx films exhibit very different properties including in- creased absorption and a metallic or semiconducting behaviour.

vii) No vacuum chamber would be required. viii) The nozzles do not clog or wear out.245

ix) The system would be easy to clean by flushing with a solvent,248 such as isopropyl alcohol.

x) Since overspraying is avoided, material consumption can be reduced by up to 80%.245 As well as avoiding waste, it was anticipated that this would reduce the build-up of

TiO2 powder in the system. 3.4 Theory of Ultrasonic Spray Deposition 67

xi) Spraying may be suitable for production if accompanied by a suitable pump, with a continuous flow, such as a gear pump or pressurized feed, and either multiple nozzles or a single traversing nozzle.46

xii) An USD system can be used to deposit a variety of films, including AR coatings and

SiO2 passivation layers. Additionally, it is possible to add other dopant liquids to the main precursor.65

◦ xiii) TiO2 films deposited by spray pyrolysis at 450 C are known to produce nearly dense, optically transparent, anatase films.177, 206

The use of an ultrasonic spray nozzle appeared to offer the best flexibility and potential for the new system. Out of the five surveyed models available on the US and Australian markets, only one (Sono-Tek Corp., U.S.A.) had been previously used in published scientific literature. This brand also had the advantage of being able to operate nozzles with a different atomisation frequency from the one ultrasonic generator.

3.4 Theory of Ultrasonic Spray Deposition

Blandenet et al. provide a good description of USD-deposited films produced by an aerosol.210 In the case of most metal oxide depositions, the aerosol is a colloidal disper- sion of a organometallic liquid in a carrier gas. The liquid is atomized in a glass vessel with a transducer operating in the kHz to MHz frequency range. The carrier gas transports the aerosol close to the heated substrate, where it is decomposed by pyrolysis. Alternately, the aerosol can react with water vapour to complete the reaction by hydrolysis. In such systems, unlike pneumatic spraying, the gas flow rate is independent of the aerosol flow rate.

In an ultrasonic nozzle, two piezoelectric transducers are contained within the nozzle, as shown in the cut-away view in Figure 3.2. This sets up a standing wave and both ends of the nozzle become anti-nodes. The junction between the two transducers is a node, a point of zero amplitude, because the transducers have their polarities opposed. This causes the transducers to either expand or contract against each other. The standing waves and anti-nodes are illustrated in Figure 3.3.248

As the TPT passes through the nozzle, the ultrasonic vibrations (48 − 120 kHz) form an aerosol. The mean diameter of the droplets produced depend upon the excitation frequency f, the surface tension σ, and the density of the liquid ρ being atomized210, 245 8πσ d = 3 k ρf 2 8πσ ∼ 3 0.34 ,(3.1) = ρf 2 68 3. TiO2 Thin Film Deposition Equipment

Figure 3.2: Cut-away view of Sono-Tek ultrasonic nozzle.248

Figure 3.3: Standing waves inside ultrasonic nozzle.248

where k is a constant, and its value of 0.34 determined by Lang.249 This relationship is plotted for a range of frequencies in Figure 3.4 for water (ρ =1g/cm3 and σ =0.073 N/m) and isopropyl alcohol (ρ =0.80 g/cm3 and σ =0.0217 N/m) at temperatures of 20−25◦C.245 The intercepts on the graph indicate the median droplet diameters for water at the excitation frequencies used in this work, 3.1µm at 48kHz and 1.7 µm at 120 kHz. Although the density of TPT is similar to that of water at 0.955 g/cm3,250 no data on the surface tension could be found. Therefore, the size of the TPT aerosol droplets could not be accurately estimated.

Although small droplet sizes are achievable with pneumatic spraying, the main advantage of USD is the narrow size distribution of the droplets.210 Figure 3.5 shows the distribution of water droplet sizes for the range of excitation frequencies 25 − 120 kHz.245 The number median diameter defines the 50% value of drop size, meaning that one half of the drops 3.4 Theory of Ultrasonic Spray Deposition 69

103

m) Water µ (

d Isopropyl Alcohol (at T=20-25oC) 102

101

Median Droplet Diameter, 100 1 10 100 1000 Excitation Frequency, f (kHz)

Figure 3.4: Median droplet diameter as a function of excitation frequency for water and isopropyl alcohol. have diameters larger than this, while the remaining half are smaller. The number mean is obtained by summing the diameters of each drop together and then dividing by the number of drops in the sample. The weight mean diameter is calculated by taking the adding together the volume of each drop in a spray sample, taking the cube root of this sum, and dividing by the number of drops. The Sauter mean diameter measures the effective ratio of drop volume to surface area and is primarily used for combustion applications.

Other factors affecting the operation of ultrasonic spraying are liquid viscosity, solids con- tent, and the miscibility of components.245 Although there are no hard-and-fast rules for determining the suitability of a liquid for ultrasonic spraying, there are several guidelines. Generally, the higher the viscosity or solids-content of a liquid, the lower the maximum flow rate. Liquids with a viscosity of up to 500 mPa s (or 50 centipoise) can be readily atom- ized,248 where water and TPT have a viscosity of 1 mPa s and 3.5 mPa s,251 respectively, at 20◦C. Liquids containing long chains can interfere with the atomization process and may inhibit the formation of discrete droplets. It has been found that liquids with a solids content of up to 40% can be successfully atomized, however the particle size should 245 be less than one-tenth of the droplet diameter. The maximum flow rate Fmax has been empirically determined to be A F = k (3.2) max f 2/3 where the constant k has the value k ≈ 28500 l Hz2/3/(s m2), and A is the atomizing surface area. Figure 3.6 plots Equation 3.2 for the range of ultrasonic nozzles produced by Sono-Tek.

The specific flow rate r is the maximum flow rate Fmax divided by the atomizing area A. The specific flow rate for the 120 kHz nozzle, not shown in Figure 3.6, is 11 l s−1m2.More importantly for research applications though, is the minimum flow rate. The minimum flow 70 3. TiO2 Thin Film Deposition Equipment

Figure 3.5: Drop size distribution for ultrasonic nozzles operating at var- ious frequencies (from Berger 245).

rate is typically about 20% of the maximum flow rate, as below this rate the liquid emerges from the nozzle in a non-uniform manner and the spray plume becomes distorted.245

3.5 TPT: The TiO2 Precursor

3.5.1 Why TPT?

Titanium isopropoxide, also known as tetraisopropyl titanate (TPT), was chosen as the

TiO2 precursor. Apart from being the most commonly used precursor in the literature, this chemical is also used on solar cell production lines. The use of any precursor can result in contamination of the TiO2 film with by-products of the precursor. Metal-organic precursors, such as TPT, often result in carbon contamination due to the residue of the organic binders.89–93, 95, 96, 98 This is typically in the order of a 1 − 2 at. % for films deposited at low temperatures. However, several researchers have observed that at higher deposition or annealing temperatures (400 − 600◦C) the carbonate species can decompose, resulting in the removal of hydrocarbon fragments.64, 99, 252 Chen et al. noted that the decomposition of TPT to TiO2 is a very clean process in which carbon is not significantly trapped either 3.5 TPT: The TiO2 Precursor 71

Figure 3.6: Specific flow rates as a function of excitation frequency (from Berger.245)

135 within the crystalline film or at the grain boundaries. Titanium tetrachloride (TiCl4)is 91, 93, 101 another common TiO2 precursor, which results in chlorine contamination. In addition corrosive by-products (HCl) are produced in the reaction.66, 91 In one instance the chlorine contamination was so high that it prevented crystallisation of the film and poor film adhesion onto the substrate resulted.101 In any case, the level of contamination observed with TPT 97 is much smaller than with TiCl4.

Further advantages of TPT are that:

i) It is non-corrosive124 and non-toxic, listed as being a mild skin and eye irritant.250

ii) It can be highly purified and has an almost indefinite shelf-.124

iii) As a liquid, it is relatively easy to handle,97 although it should not be exposed to a naked flame.250 The fact that it is not dangerous makes the addition of TPT to a CVD system a relatively easy and safe task, as no special gas handling equipment is required.97

iv) It is very volatile at low temperatures (50◦C), which means that it will be readily decomposed.97

v) It can be ultrasonically sprayed directly without dilution.49

vi) It has been observed that there is enough oxygen in the TPT that the reaction 94, 99, 101, 124, 221 to form TiO2 can proceed without additional oxygen in the ambient. 72 3. TiO2 Thin Film Deposition Equipment

3.5.2 The TPT→TiO2 Reaction

The mechanism of the reaction of TPT aerosol to form TiO2 depends on the droplet size. If the majority of the aerosol is a gas when it contacts the substrate then the deposition conditions are similar to a CVD process.210 Aerosols with larger droplet sizes will not have had time to vaporise completely, while smaller droplets will be decomposed by pyrolysis be- fore striking the substrate. Since the droplet size distribution is small in ultrasonic spraying, the same decomposition conditions will apply to nearly all of the aerosol. The deposition conditions vary also with temperature, as the diagram for pyrolysis in Figure 3.7 shows, however other factors such as flow rates and geometry also play a role. In process A, the decomposition rate at very low temperatures (< 100◦C) will be slower than the deposition rate and a liquid film will form on the surface. This layer will slowly dry, however it will still contain many organics and probably cracks.115 In process B, the droplets evaporate before reaching the surface and a precipitate strikes the substrate where decomposition occurs. In process C, the solid precipitate melts and vaporises (or sublimes) and the vapour diffuses to the substrate and undergoes a reaction there. This corresponds to true CVD. At higher temperatures (process D), the vapour undergoes a chemical reaction before impinging upon the substrate. The droplets in the aerosol have formed solid particles that stick to the surface of the substrate. The product information sheet on DuPont’s “TYZOR” TPT notes that TPT pyrolyses at temperatures greater than 350◦C, and that films deposited in this method at 500 − 600◦C are considerably harder than films produced by hydrolysis and contain no organic residue.252 VXEVWUDWH ILQHO\GLYLGHG VROLGSURGXFW

YDSRXU

SUHFLSLWDWH

{ { { { GURSOHWV

$ % & '

/RZWHPSHUDWXUH +LJKWHPSHUDWXUH

Figure 3.7: Pyrolysis decomposition as a function of temperature (adapted from Vigui`eand Spitz 244).

253 The decomposition of TPT (by pyrolysis) to form TiO2 proceeds as follows:

Ti(OC3H7)4 −→ TiO2 +2C3H7(OH) + olefins , (3.3) 3.5 TPT: The TiO2 Precursor 73

For the reaction of TPT to form TiO2 by hydrolysis, the reaction product will be strongly dependent on the amount of water vapour present in the system, as well as the substrate temperature. Wong et al. found that for APCVD-deposited TiO2, there is not enough oxygen to form TiO2 and the TPT remains unreacted on the wafer when less than 30% relative humidity exists.204 The reason why the TPT did not react by pyrolysis in this ◦ case (Tdep =250C) remains unclear. When spraying in an environment with greater than 45% relative humidity, the TPT reacts before long before reaching the wafer and a powdery white deposit results.46, 202, 203 Wong et al. mentioned that laboratory scale CVD experiments showed it was possible to deposit films with 15% thickness uniformity at humidities above 45%.204 In general, the recommended range in relative humidity is from 30% to 45%, which results in a transparent homogeneous film being deposited.202, 203 In this case, the reaction rate is limited by the amount of TPT reaching the wafer,205 which decomposes by hydrolysis onto the heated wafer, according to the two-step hydrolysis/degradation reaction in Equation 3.4.210, 253 The “Tyzor” TPT product information sheet (Du Pont, Inc.) contains a complete explanation of all the steps involved in the hydrolysis reaction.252 In the information it is also noted that whether hydrous titanium dioxide (TiO2·H2O) or TiO2 itself is formed as the final product is dependent on the temperature and the rate at which water is added to the system.

Ti(OC3H7)4 +2H2O −→ Ti(OH)4 +4C3H7(OH)

Ti(OH)4 +4C3H7(OH) −→ TiO2 +4CH3CH(OH)CH3 (3.4)

In Equation 3.4, one mole of TPT reacts with two moles of water vapour to form one mole of TiO2 and 4 moles of by-product (mostly 2-propanol). Practically however, the amount of water vapour will depend on the geometry of the reaction zone, and gas flow and exhaust 217 rates. TiO2 films deposited in an APCVD system required more than four times this amount of water vapour for the above reaction to occur.204 As the substrate temperature is increased, the deposition chamber, or environment, will heat up and the relative humidity will decrease. However this is not strictly of great concern, as the number of moles of water vapour going into the system will remain constant, and this is what is required.

The amount of TPT and water vapour required for the reaction to occur, as given in Equation

3.4, can be calculated. The mm of TiO2 is 79.9 g/mol. The number of moles M 2 −6 3 of TiO2 to cover a 25 cm surface area with a 70 nm thick film (volume, ν =175× 10 cm ) is given by the following equation, where the density ρ of rutile is 4.26 g/cm3 ρν M = mm −6 =9.3 × 10 moles TiO2 (3.5)

−6 −6 This means we require 9.3 × 10 mol of TPT and 18.6 × 10 mol of H2O for the reac- tion in Equation 3.4 to occur. The molar mass of TPT is 284.26 g/mol and its density is 74 3. TiO2 Thin Film Deposition Equipment

0.955 g/cm3.250 Rearranging Equation 3.5 we obtain Mm ν = m ρ =2.8 µl of TPT to cover the 25 cm2 area with a 70 nm thick film. (3.6)

The amount of water vapour required has been estimated at being four times greater again due to unreacted water vapour being extracted out the exhaust.204 However, in the system constructed for this work, many times this amount (approximately 0.5 − 1 ml) of TPT is needed to be sprayed as the film density is lower than that of single crystal rutile and not all the TPT sprayed is deposited within this area. This is especially true for shallow deposition angles, where the majority of the TPT passes across the surface of the wafer. Based on the requirement of 1 ml of TPT being required for each 5 cm × 5 cm wafer we can expect that at least 4 ml of water vapour would be required.

3.6 Design of Ultrasonic Spray Deposition System

3.6.1 Selection of Ultrasonic Nozzle

Two ultrasonic atomizing nozzles (UAN) were purchased from Sono-Tek Corporation (Mil- ton, N.Y., U.S.A.). There were several reasons for selecting an ultrasonic nozzle, rather than an ultrasonic nebulizer. Firstly, with a nozzle the spray could be directed at different angles. Secondly, it was believed that a spray system where the reaction mechanism could be altered with the drop size could be tuned to behave similarly to a belt-furnace APCVD system204 as used in the PV industry. Thirdly, the potential of a solely spray-based processing system for fabricating solar cells was attractive. As well AR coating deposition, spray systems have been used for depositing passivation layers for solar cells.208

The two nozzles operated at frequencies of 48 kHz and 120 kHz, respectively. One motivating factor for nozzles manufactured by Sono-Tek was that only one broadband ultrasonic gener- ator (BUG) was required to operate any of their nozzles. In contrast, other manufacturers require that a separate BUG be purchased for each nozzle. The 48 kHz nozzle was selected in order to evaluate the system performance at high flow rates, and also to determine whether large droplets would be prevented from entering the 20 − 30 µm wide grooves on the front surface of the wafer. The 120 kHz nozzle was selected because it was able to tolerate the lowest flow rates. Later on in the project the 48 kHz nozzle was exchanged for a newly released nozzle called the MicroSpray. This nozzle used a 120 kHz atomizing frequency as well, however it was fitted with a micro-bore tube which significantly reduced the inside diameter of the nozzle. This enabled extremely low flow rates to be tolerated, while still achieving acceptable spray plumes. Figure 3.8 illustrates the 48 kHz standard nozzle and the 120 kHz MicroSpray nozzle purchased from Sono-Tek. The 120 kHz standard nozzle is 3.6 Design of Ultrasonic Spray Deposition System 75 similar in appearance to the MicroSpray nozzle, except that it possessed a conical tip (see Figure 3.8).

Figure 3.8: The 48 kHz (left) and 120 kHz (right) ultrasonic nozzles pur- chased from Sono-Tek Corp (adapted from Sono-Tek Corp.246).

The front and rear horns of the nozzles are manufactured from , Ti-6Al-4V, while the nozzle housing and liquid inlet are made from 316 stainless steel. The titanium alloy is chosen for is high mechanical strength, good acoustical properties and excellent chemical resistance.245 Most chemicals, except hydrofluoric and sulphuric acid and strong oxidizing agents, are compatible with these spray nozzles. Table 3.1 provides technical details on the Sono-Tek MicroSpray ultrasonic atomizing nozzle as well as the standard 120 kHz and 48 kHz nozzles. The dimensions in Table 3.1 correspond to the schematic of the nozzle shown in Figure 3.9.

Figure 3.9: Dimensions of the Sono-Tek ultrasonic atomizing nozzles that correspond to the values given in Table 3.1 (adapted from Sono-Tek Corp.246). 76 3. TiO2 Thin Film Deposition Equipment

Table 3.1: Sono-Tek ultrasonic atomization nozzle data (from Sono-Tek Corp.246). Note that dimensions are not exact as they are converted from inches. So few experiments were performed with the 48 kHz nozzle that some parameters are unknown and are labelled in the table as not applicable (n/a).

Nozzle Type Microspray UAN UAN Parameter UAN Atomisation frequency (kHz) 120 120 48 Orifice diameter (mm) 0.38 1.32 2.18 Maximum flow rate (ml/min) 2.4 21 72 Recommended minimum flow rate (ml/min) 0.48 4.2 14.4 Actual minimum flow rate (ml/min) 0.02 0.35 n/a Atomization power at minimum flow rate (W) 1.1 − 2.0 1.1 − 2.0 n/a Median drop diameter (µm) 18 18 38 Weight (g) 196 309 Dimensions: A1 (mm) − 5.8 11.7 B1 (mm) − 11.2 26.9 A2 (mm) 2.5 − − B3 (mm) 11.5 − − C (mm) 29.3 25.4 37.3 D (mm) 36.6 36.6 38.1 E (mm) 12.7 12.7 42.9 F (mm) 10 8.6 −

3.6.2 Ultrasonic Nozzle Performance

It was quickly determined that the 48 kHz nozzle would be unsuitable for our application. This was due to the extremely high flow rates, which caused a much thicker film than desired to be deposited in less than one second and was difficult to control. The 120 kHz proved to be better with spraying times in the order of 1 min for a 70 nm thick film at the lowest flow rates. Both of these nozzles possessed a conical tip, as shown in the left hand side of Figure 3.9, which is designed to spread the spray plume out over several inches. This tip was initially favoured as it was believed that this coverage would be sufficient to coat a 2” wafer. However, to achieve longer deposition times, the lowest flow rates had to be used. Although a fine mist often emerged from the nozzle the low volumes being pumped were not sufficient to create a plume of several inches in diameter. Occasionally, the nozzle would also “stall” and spit larger droplets of liquid onto the substrate. This resulted in 1 − 30 µm diameter TiO2 particulates being incorporated into the film. These defects reduced the chemical resistance of the films.

To improve the film uniformity, reduce particulate incorporation, and to achieve lower flow 3.6 Design of Ultrasonic Spray Deposition System 77 and film deposition rates the 48 kHz nozzle was exchanged for a 120 kHz MicroSpray nozzle. This nozzle had a much narrower bore enabling very low flow rates could be used (down to 0.02 ml/min). Additionally, the tapered tip (see tip on right-hand side of Figure 3.9) resulted in a much finer line of spray. Achieving a consistent spray was aided by the use of the small-bore 1 ml gas-tight syringes (see Section 3.6.3). This setup resulted in a deposition time of 1 − 2 min and the best performance of all configurations tried for the spray system. The power required from the broadband ultrasonic generator (BUG) to atomize the TPT with the 120 kHz nozzles was in the order of 1.5W.

3.6.3 Liquid Delivery

Pump

Selecting the correct pump is a crucial step in designing the spray deposition system. This is because the ultrasonic nozzle will atomize any liquid that reaches the atomizing surface. A syringe pump (Yale YA-12, Kent Scientific, U.S.A.) was chosen primarily because it can pump continuously in the µl/min range and pulsed volumes of nanolitres can be achieved.245 The upper limit on the pumping rate and the dose is limited by the maximum syringe size, in this case 60 ml. The pumping action is very smooth, which is important for achieving good spraying. The pump is programmable and can be controlled via TTL signals or an RS-232 port. In this work it was controlled via the front panel only. This enabled a flow rate, typically 0.02 − 1.00 ml, to be set and the pumped volume to be monitored on the display. It is also possible to inject or withdraw with the pump. The ability to withdraw liquid was used in conjunction with a three-way valve in order to refill the smallest (1 ml) syringes (refer Figure 3.11).

Figure 3.10: The Yale YA-12 syringe pump. 78 3. TiO2 Thin Film Deposition Equipment

Syringes, Tubing, Valves and Bottles

Initially experiments were performed with disposable 5 ml and 10 ml plastic syringes with rubber-tipped plungers. By observing the spray plume it was noticed that the plunger was periodically sticking as it travelled down the syringe barrel. It was discovered that the ability of a syringe pump to dose at consistently low flow rates is influenced by the diameter of the syringe. Therefore, several high-quality 1 ml and 2.5 ml gas-tight syringes equipped with Teflon plungers were purchased from a chromatography supplier (SGE, Melbourne, Australia). These performed excellently and could be flushed with dilute HF to remove any build-up of TiO2.

All the syringes used were of the Luer-lock variety to ensure an air-tight and leak-proof connection. A variety of chemically resistant plastic (HDPE) fittings were purchased (Chro- malytic Technology, Victoria) to adapt from Luer-lock to barbed fittings for 1/8” outer diameter (1/16” inner diameter) tubing. The 1/8” plastic tubing (LDPE) then connected directly to the 1/8” Swagelok liquid inlet fitting on the Sono-Tek nozzle.

As mentioned previously, a Teflon-lined gas-tight three-way valve (SGE, Australia) was later installed to permit withdrawal of new TPT precursor from a 100 ml screw-top chemical bottle. This was convenient for the 1 ml syringes since coating a 2”diameter wafer with a

70 nm thick TiO2 typically required 0.5 ml of TPT. The gas-tight syringes were able to screw directly into the valve, however it was necessary to purchase Kalrez fittings (SGE, Australia) to connect the 1/8” tubing.

Figure 3.11 indicates how the Schott chemical bottles were modified to contain TPT. Initially a hole was drilled in plastic lid and a thread tapped into the hole. A female Luer adapter was screwed into the lid, sealed with a Viton o-ring, and the 1/8” tubing extended down into the bottle. A Teflon solvent filter (SGE, Australia) was pressed onto the end of the tubing, and this rested about 5 mm above the bottom of the bottle. This was to remove as many particulates as possible from the TPT. Particulates can slowly form due to the bottle being opened and closed, and also if it is not totally air-tight.

3.6.4 Substrate Heater

The previous substrate heater for the pneumatic spray system at UNSW consisted of a stain- less steel block that was placed on a 2 kW stove element.203 This was less than ideal as the majority of heat generated did not pass into the block. Therefore, a new block was designed. Stainless steel was still used as it was anticipated that the block would need to withstand temperatures of 500◦C or more. Other materials considered included copper, graphite, nickel, and titanium. Copper is an undesirable material for semiconductor devices due to the speed that which it can diffuse through silicon. The remainder of the materials were 3.6 Design of Ultrasonic Spray Deposition System 79

2' (QG /'3(WXELQJ FDS 0DOH/XHUWR EDUEDGDSWHU

)HPDOH/XHUILWWLQJ VFUHZHGLQWROLGZLWK 9LWRQ IOXRULQDWHG UXEEHU RULQJ

2' /'3(WXELQJ

6ROYHQW )LOWHU

Figure 3.11: The creation of a filtered TPT reservoir from a chemical bottle.

ruled out because of expense. Therefore, a 150 mm×150 mm×20 mm stainless steel block was milled out for use as a substrate heater. Figure 3.12 shows the workshop drawings of the block, including precisely milled holes for the six heater cartridges, thermocouple and vacuum line. Heater cartridges were selected as they would result in the best power to heat conversion. The heater cartridges chosen were a split design (Dalton Electric Co., MA, U.S.A.) that expand upon heating to ensure good thermal contact to the block (see Figure 3.13). This avoided the use of a copper thermally-conducting paste. The power rating of each cartridge was 250 W and could be run off mains power (240 VAC . The cartridges were custom designed so that a 20 mm length near the leads was not embedded into the block. This cool-zone was important for keeping the shielding around the leads from burning or melting and possibly shorting out.

Double layer fibre-glass/ceramic sheets (6 mm thick each) were placed on five sides of the block, and were contained by a stainless steel box with 1 mm thick walls. The insulation was designed to prevent too much heat from being transmitted downwards, as this could warp the rails of the driving mechanism below. A 3 mm diameter K-type thermocouple was inserted into the block. The tip of the thermocouple was located underneath the centre of the wafer. The output from the thermocouple was fed into a BTC-9090 PID temperature controller. The output of the PID controller to the six heater cartridges was via a 240 VAC 10 A relay. As the heater cartridges could draw a maximum of 1.5 kW there was no risk of 80 3. TiO2 Thin Film Deposition Equipment

3 BACK Hole for thermocouple 3mm 20 VIEW diameter, 70mm deep 150 5 75 3 45 35 25 15 Vacuum 3mm diam. hole channels: for thermocouple 1mm wide, 75 1mm deep

150 TOP VIEW

5

3mm diam. vacuum line

1/8” NPT 10 thread 50 25 6x holes for heater 25 cartridge = 0.377” 50 50 3 10 FRONT 0.377” VIEW 20

Figure 3.12: Schematic diagram of the stainless steel heater block. 3.6 Design of Ultrasonic Spray Deposition System 81

Figure 3.13: Six of these 6” long 250 W Dalton Watt-Flex heater car- tridges were embedded in the stainless steel block to ensure maximum heat transfer.254 blowing the fuse on the output side of PID controller. The block took about 15 min to reach 450◦C from room temperature, and typically temperature fluctuations while spraying were not more than 5◦C.

3.6.5 Motorized Stage

A motorized translation stage was designed so that the substrate could pass back-and-forth beneath the spraying zone. The 12 VDC motor (RS Components, Sydney) had a continuous torque rating of 300 mNm and a maximum torque rating of 600 mNm. At 12 VDC the motor speed was 220 rpm. The motor was directly connected to a double-start threaded spindle with a 2 mm thread pitch. This spindle passed through a Teflon nut that was mounted on a plate connected to the underside of the heater block. The significant mass of the heater block (about 6 kg) was supported by two 12 mm diameter stainless steel rails. These 50 cm long rails were machined precisely to accommodate four linear bearings. Once correctly aligned, the block slid effortlessly along the rails. Thus, by applying either a positive or negative voltage to the motor the spindle turned and moved the block forwards or backwards. At 220 rpm the block travelled at a speed of 7 mm/s.

As it was necessary to keep the door to the spray chamber shut during depositions (see

Section 3.6.7), a simple circuit using a double-pole change-over (DPCO) 12 VDC relay was implemented in order to reverse the direction of the stage. The relay has two pairs of inputs. The first pair consisted of +12 VDC and -12 VDC, while the second had the polarity switched. Two sealed and chemically resistant momentary push-button micro-switches (RS Components, Sydney) were used to trigger the state of the relay. Pressing one switch caused the relay to “change-over” from its first input to its second input pair, while pressing the second switch caused the relay to go back to its initial state. Thus, the stage started moving in one direction (the direction it was last travelling) until reaching the end of the rails upon which the microswitch was depressed, and the stage changed direction. A fuse was also included in this simple circuit to prevent the motor from burning out. A motor speed controller was also added to the system to enable slower translation of the stage, however this was rarely used. A 2 A 12 VDC power supply was sufficient to power all the electrics. Finally, in order to keep the rails free from TiO2 dust, concertina-style rubber boots were fitted over the rails. These boots could be compressed to about 25% of their standard length. 82 3. TiO2 Thin Film Deposition Equipment

3.6.6 Spray Shaping

Spray shaping and directional control was necessary for several reasons. In early experiments it was found that when the nozzle was placed directly above the substrate heater the fine, low-velocity (about 10 cm/s) mist did not fall on the wafer. This was because of hot air rising from the 450◦C block and carrying the spray droplets away with it. Experiments were performed with the two “air shrouds” that were supplied with the nozzles (Sono-Tek Corp.). This enabled the spray to reach the heated substrate, however there were many particulates in the film and the thickness uniformity was very poor.

After discussing our requirements with Sono-Tek Corp., a loan of a “vertical spray assembly” was arranged. The vertical spray assembly, depicted in Figure 3.14, is designed to produce wide spray patterns. The two streams of gas are slightly off-centre, resulting in the shearing of spray plume.245 Wide spray patterns were able to be generated with the vertical spray assembly, however these were far from uniform. Additionally, whatever nozzle-substrate distance was used resulted in the incorporation of many particulates into the film. It is believed that the TiO2 particulates arise from droplets reacting with remnant humidity in the air, entrained into the nitrogen gas stream. Since the nozzle is directly above the wafer any particulates formed will fall onto the wafer and be incorporated into the film.

As a solution, an “air-knife”, sometimes also called an air-guide, was purchased (Exair Corp., U.S.A.). The air-knife was 6” long and consists of two halves of an aluminium casing bolted together. There is an 1/4” NPT fitting at one end for the nitrogen inlet. At the front there is a small slit running almost the whole length of the air-knife, which creates a high-velocity sheet or curtain of nitrogen. The slit height is set by a plastic shim inside the air-knife. It was found that by intersecting the emerging spray plume from the nozzle with a sheet of nitrogen from the air-knife enabled the deposition of a visually acceptable film on the wafer. As the 6” wide curtain cooled the substrate significantly a new shim was inserted, which had created a slit with 4 mm width and 0.02” height. This reduced the amount of cooling of the block significantly.

3.6.7 Miscellaneous Equipment

Relative Humidity Sensor

As discussed previously in Section 3.5.2, the reaction of TPT to form TiO2 is extremely sensitive to the relative humidity. Relative humidity RH is defined as the partial pressure p of water vapour in air divided by the vapour pressure of water ps at the same temperature. To be able to monitor the relative humidity in the spray system, a sensor and digital meter were purchased (Elan Technical Corp., CT, U.S.A.). 3.6 Design of Ultrasonic Spray Deposition System 83

Figure 3.14: Sono-Tek’s vertical spray assembly, designed for generating wide spray patterns (from Sono-Tek Corp.246).

Oxygen Concentration Sensor

An was added to the system because it was anticipated that spray depositions with TPT diluted in a solvent, such as isopropanol, would be performed. Although the volumes of solvents sprayed would be less than 1 ml, ensuring that the oxygen concentration was low would inhibit the combustion of the solvent. Therefore an oxygen sensor (Electrovac GmbH, Austria) was added to the system. As shown in Figure 3.15(a), when a voltage is applied across the zirconia electrolyte cell, oxygen is pumped through the cell because oxygen ions carry the current through the cell. By attaching a cap with a pinhole on the cathode side of the cell and increasing the voltage, the current becomes saturated due to limited transfer of oxygen ions to the cathode. This current is proportional to the ambient oxygen concentration.255 The advantages of this sensor included having a linear output signal, no cross-sensitivities to other gases, long life, and a very low temperature dependence of the 84 3. TiO2 Thin Film Deposition Equipment signal. Due to a heater inside the package the sensor required the application of 2 V while warming up (30 s) and then 4 V for constant operation. Figure 3.15(b) shows an image of the oxygen sensor. Designing this power supply and installation of the sensor and a suitable analogue meter was completed by a German practicum student, Manfred Fahr.

Figure 3.15: (a) Schematic indicating oxygen sensor operation, and (b) Image of the oxygen sensor (adapted from Electrovac GmbH 255).

Spray Chamber

The spray chamber was fabricated from steel “speed-frame” and clear perspex sheeting, with internal dimensions of 70 ×70 ×70 cm. The entire front side of the chamber was hinged along one edge and could be opened to provide full access to the chamber. A series of 25 mm diameter holes were drilled in the door and covered with a sliding plate. This was in order to allow the entry of ambient air after spraying was completed. The top of the chamber was connected to a 6” diameter exhaust line. The exhaust flow could be controlled via a large butterfly valve.

Nitrogen Heater

Due to the relatively high gas flow rates used (see Section 3.6.8) significant cooling of the block occurred. With the block temperature set at 450◦C the top surface of the wafer was measured to be about 330◦C while spraying. In order to reduce the amount of cooling occurring it was decided to control the temperature of the nitrogen gas stream. The nitrogen heater used a heater cartridge that was inserted inside a length of 3/8” stainless steel tubing. It was designed to operate upright so that if the nitrogen gas flow stopped the hot air for around the heater would rise up to the top of the tube where a thermocouple was placed. The system was designed to operate at temperatures up to 400◦C, being controlled by a BTC-2020 PID controller (ECE Fast, Melbourne). 3.6 Design of Ultrasonic Spray Deposition System 85

Air-Knife Heater

Although the nitrogen heater worked well, a with heated lines the temperature of the emerg- ing gas had dropped only 30◦C after passing through a 1 m length of tubing (with the heater set to 200◦C). However, when connected to the aluminium air knife, which acted as a very efficient heat sink, the emerging gas temperature was only a few degrees above ambient temperature. Therefore a 240 VAC 100 W heater pad was obtained (RS Components, Aus- tralia) that was both flexible and had a high-temperature adhesive applied to one side. This was carefully folded around the air-knife and firmly held in place. This permitted heating (not-controllable) of the air-knife up to temperatures of 190◦C, which enabled the heated nitrogen gas to reach the substrate.

3.6.8 Operation of the TiO2 Spray System

In its final form, the spray system incorporated all of the above componentry, as well as regulators and flow meters to accurately control gas flows (as shown in Figure 3.1). All the films deposited using this deposition system were formed by spray pyrolysis, due to the very low relative humidity. The diagram in Figure 3.16 indicates the necessary steps to set-up and perform depositions with the TiO2 spray system.

Table 3.2 in Section 3.7 lists the various deposition parameters of the system. The gas flow rates listed were optimised in order to extend the film coverage in the forward direction, while trying to reduce the amount of cooling of the substrate heater. At lower flow rates, the film was not deposited onto the wafer due to heat arising from the 450◦C block, while little benefit was achieved by using higher flow rates. A typical atomisation power of 1.5 − 2.0W was required from the BUG for flow rates between 0.02 − 0.10 ml/min. Spray depositions were only performed at the maximum system operating temperature of 450◦Casatlower temperatures the number of particulates in the film was unacceptable. It should be noted that the deposition temperatures quoted throughout this work may be slightly higher than the actual deposition temperature due to the cooling of the wafer by the gas flow.

The deposition time and the efficiency at which the TPT is converted to TiO2 is greatly influenced by the alignment of the system. The height of the nozzle and whether the tip sits directly in the nitrogen flow from the air-knife or just above it is important. The best results were achieved by adjusting the height of the nozzle with the nitrogen flowing, and when it could be heard that the tip just entered the gas stream the nozzle was fixed in that position. The optimum angle for spraying was determined to be 5◦ below the horizontal. At this angle the nitrogen emerging from the air-knife travels up towards the block and then adheres to the top surface of the block as it passes across it. The adherence of a gas to a solid surface is called the Coanda effect. In general, spraying at angles closer to the horizontal reduced the number of particulates observed in the film, and also enabled only the top surface of 86 3. TiO2 Thin Film Deposition Equipment

Check that tip of spray nozzle is free TiO2 powder

Install a new 0.5 m length of LDPE 1/8" OD tubing between three-way valve and nozzle

Withdraw 1 ml of TPT into gas-tight syringe

Syringe pump: set syringe diameter and desired flow rate

Purge chamber with N2 until RH ≈10%

Wait for substrate heater to reach 450°C

Inject TPT to fill "dead-space" in the tubing, valves and nozzle

Turn on motor to move stage away from centre and place wafer on substrate heater

Using three-way valve, withdraw TPT from bottle into syringe

To begin spraying: adjust power on BUG to 1.2 W, start TPT flow, and turn on air-knife N2

Leave motor off momentarily as there are often a few "spits" more wafers at the start of spraying y ra p

Turn on motor and monitor colour of TiO2 film on wafer - stop s To when it appears dark blue

Turn off syringe pump, motor and BUG, and remove wafer. Open exhaust valve and purge system for 1 min.

To finish spraying, turn off heater block, remove nozzle and syringe and rinse with IPA, discard tubing in waste container

Figure 3.16: Diagram indicating the necessary steps to set-up and per- form depositions with the TiO2 spray system. 3.7 Design of CVD System 87 groovedwaferstobecoated.

The order in which the various components are turned on and off are important. The tubing, valve and nozzle should be filled with TPT before starting. Any excess TPT can be wiped off the tip of the nozzle using a tissue and some isopropyl alcohol (IPA). The stage should be positioned so that the initial spray plume will not land on the wafer. This is because of a tendency for there to be more particulates in the initial spray plume. Secondly, the air-knife flow should be started and the BUG should be turned on. After a few seconds of spraying the motor can be switched on and coating of the wafer will begin. To finish, it is best to turn off the BUG first, to stop he atomisation process as quickly as possible, then the syringe pump, and finally the motor and air-knife flow.

For the deposition of doped TiO2 films (as in Section 6.3 of this work) the dopant liquid was mixed together with the TPT in a small chemical bottle in a fumecupboard before spraying. If many doped films, especially of different doping concentrations, were required it would be recommended to replace the three-way valve with a four-way valve and to have separate pure bottles of TPT and the dopant liquid. The syringe should be shaken to ensure good mixing of the two liquids.

3.7 Design of CVD System

3.7.1 Motivation

Although the spray pyrolysis system produced quite dense anatase films, the thickness uni- formity on a macroscopic level was poor. On a scale of nanometres to micrometers the films were quite uniform, however, on a scale of millimetres to centimetres, thickness variations of a similar magnitude to the average film thickness were observed. This meant that, for a 70 nm thick film, there were many points across the wafer that remained virtually uncoated. This, along with the incorporation of particulates, typically 1 − 30 µm in diameter, had se- rious implications for chemical resistance and the electroless metal plating process. Other limitations included only being able to operate the spray deposition system at the maximum temperature of 450◦C, in order to limit the number of particulates in the film.

3.7.2 TPT Bubbler and Temperature Control

Therefore a simple CVD system was designed to replace the spray system. Attachai Uer- anatusan (MEngSc, UNSW) had performed initial investigations into a CVD nozzle that was fed with a vapour from a glass bubbler. This idea was adapted by the author into the existing structure of the spray system. A one-litre stainless steel bubbler was purchased (Meriter, U.S.A.) for safety as it could handle pressures up to 42 psi, much greater than the 88 3. TiO2 Thin Film Deposition Equipment quartz bubbler. The use of electropolished stainless steel bubblers have been successfully demonstrated in the semiconductor industry.256 It was found that there was essentially no assay degradation when comparing the performance of quartz and stainless steel bubblers (J.C. Schumacher, CA, U.S.A.) were the same and as long as moisture was prevented from entering the bubbler. This is achieved by using metal gaskets instead of Teflon. A pressure relief valve with a threshold of 35 psi was placed in parallel with the bubbler to prevent too high a pressure being applied to the bubbler. The outlet of the pressure relief valve went directly into the exhaust system.

A Schumacher temperature controller and temperature control unit were used to maintain the bubbler temperature at 50◦C. At this temperature, the TPT has a vapour pressure of 1 mbar. The Teflon tubing leading from the bubbler to the nozzle was heated using a

1 m length of 240 VAC heater tape with a power of 90 W. This was insulated with glass- fibre insulation tape, held in place with high-temperature heat-shrink tubing. This was to prevent condensation of the TPT onto any cool surface so as to avoid the build up of TiO2 particulates in the tubing and nozzle. It was not necessary to hear the nozzle itself as its close proximity to the heater block ensured that it maintained a temperature well above 50◦C.

3.7.3 Water Vapour Bubbler

After initial success with TiO2 films deposited using the CVD system, a second bubbler containing de-ionised (DI) water was added to the system. This was designed to permit reactions by hydrolysis. This was a standard quartz bubbler typically used for performing wet oxidations in a tube furnace. A separate nitrogen regulator was used to limit the pressure that could be applied to the bubbler to 5 psi. The bubbler had a matching base with an integrated heater. The temperature was monitored with a thermometer and manually maintained at approximately 100◦C. A 3-way valve was attached to the outlet of the water bubbler so that the water vapour could be ”switched-off” by redirecting the flow to an empty flask. Figure 3.17 shows a diagram of the final CVD system.

3.7.4 Operation of the CVD System

Operation of the TiO2 CVD system was quite similar to the spray system. About three hours before depositions were performed it was necessary to switch on the temperature controllers for the TPT and water bubblers. A timer was implemented for doing this so that CVD depositions could begin early in the morning. Depositions could now be performed at substrate temperatures of 150 − 450◦C. Below 150◦C the film coverage became very non- uniform, probably due to the TPT striking the substrate as a liquid. As with the spray system the air-knife angle and nozzle position were critical. It was found that the optimum 3.8 Conclusions 89 angle for the air-knife was about 2◦ above the horizontal, and the height on the nozzle was adjusted in the same manner as previously described in Section 3.6.8. Again, the gas flows listed in Table 3.2 are the minimum flows for successful TiO2 deposition to occur, while trying to reduce the amount of cooling of the block occurring.

Table 3.2: Deposition conditions for the USD and CVD TiO2 films Process parameters USD CVD Liquid TPT flow rate (cm3) 0.02 − 0.10 n/a TPT bubbler flow rate (scm3) n/a 1700 Air-knife N2 flow rate (slpm) 5 13 Atomisation power (W) 1.5 − 2.0 n/a Chamber pressure (kPa) ≈ 101 ≈ 101 TPT temperature (◦C) 25 50 Substrate temperature (◦C) 450 150 − 450 Nozzle-substrate distance (cm) 5 − 7 3 − 5 Relative humidity (%) < 10 < 10 Deposition angle (◦) −5 ≈ 2 Deposition time (min) 1 − 2 10 − 20

3.8 Conclusions

A good understanding of basic TiO2 material properties along with knowledge of a previous

TiO2 spray deposition system at UNSW, formed the requirements for the TiO2 thin films in this application. Films were required to be dense and defect-free, exhibit good thickness uniformity, possess a high refractive index and low optical absorption, and be insulating. Examination of deposition techniques described in the literature lead to the author design- ing and constructing an ultrasonic spray deposition (USD) system. USD offered several many potential advantages, including shallow-angle depositions, low deposition rates, and the ability to deposit for a wide range of thin films via this method using different liquid precursors. Many of the desired TiO2 film properties were obtained from films deposited using the USD system. However, the limitations of the USD system were, firstly, that the thickness uniformity of these films was poor, secondly, that large particulates were commonly embedded into the thin films (these films will be characterised in detail in Chapter 4). Spray depositions were performed at 450◦C as a dramatic increase in the number of particulates was observed at lower temperatures. Thus, although the films exhibited a high refractive index, the ability to tune the refractive index by varying the substrate temperature could not be realised. This lead to the design and construction of a simple chemical vapour deposition

(CVD) system that could operate at atmospheric pressure. TiO2 films deposited with this simple CVD system are used throughout the majority of this work. 90 3. TiO2 Thin Film Deposition Equipment PID contoller temperature contoller temperature TPT bubbler C C C C q q q q 445 450 49.9 50.0 DC 12 V 2 N line

switch (x2) switch thermo -couple vacuum st. steel block relay housing st. steel bubbler TPT heater rings vacuum motor + 1/8" st. 1/8" steel nozzles 2 TPT N heaters lines cartridge heated O + 2 2 H N 2 N O 2 r heate quartz H quartz air-knife bubbler and bubbler 2 N

Figure 3.17: TiO2 CVD deposition system. The humidity sensor, exhaust valve on TPT bubbler, and the 3-way valve on the water bubbler have been omitted for clarity. Chapter 4

Characterisation of TiO2 Thin Films

Extensive characterisation of TiO2 films deposited using ultrasonic spray deposition (USD) and chemical vapour deposition (CVD) was performed in order to determine the physical, optical, electrical and chemical properties of the films. All films deposited at 450◦C were of the anatase phase. A surprising result was that USD anatase films did not convert to rutile ◦ 3 after lengthy annealing at 950 C. USD TiO2 films were found to be dense (3.64 g/cm ) and exhibited a high refractive index (2.45 at 600 nm), ideal for acting as an antireflection (AR) coating on a glass encapsulated silicon wafer. Although the USD films appeared continuous over a microscopic level, macroscopically the films exhibited large variations in thickness and the occasional pinhole. TiO2 films deposited using CVD were found to exhibit much lower surface roughness and better thickness uniformity, although the density was significantly lower than that of the USD-deposited films. Shallow-angle depositions were successful in maintaining the grooves free of TiO2. However, shallow-angle depositions were not successful on textured crystalline silicon wafers, with large tree-like structures growing at the tips of the pyramids. The chemical resistance of all films was excellent against acids, but relatively poor against alkaline solutions, although the etch resistance did improve upon annealing of the

TiO2 film.

4.1 Introduction

The literature review in Chapter 2 described the many different material phases of TiO2, and the variation of the material properties with those phases. The front-surface dielectric film in the buried-contact (BC) solar cell fabrication sequence is subjected to a number of high-temperature processing steps in oxygen, nitrogen, and phosphorus-containing furnace ambients. Therefore, before replacing the silicon dioxide (SiO2) layer in the original BC solar cell with a TiO2 film, it was necessary to perform extensive characterisation of the TiO2 thin

films to, firstly, optimise the deposition parameters to obtain a TiO2 film exhibiting the

91 92 4. Characterisation of TiO2 Thin Films desired qualities and, secondly, to determine the behaviour of the films under typical BC fabrication sequence conditions. Different characterisation techniques were utilised in order to determine the physical, optical, chemical and electrical properties of the as-deposited and annealed TiO2 thin films, deposited using ultrasonic spray deposition and chemical vapour deposition (CVD). Characterisation techniques can be roughly divided into the following four categories:

• Physical Properties: The crystalline phase of the film and the existence of oxygen vacancies were determined using Fourier-transform infrared (FTIR) and Raman spec- troscopy. The surface roughness of the films was measured using atomic force mi- croscopy (AFM). The presence of particulates and the density of the films was ob- served using scanning electron microscopy (SEM). SEM work also provided a rough indication of the grain size of the material. Elemental analysis of film was performed using both X-ray photoelectron spectroscopy (XPS) and Rutherford back-scattering (RBS) spectroscopy.

• Optical Properties: The refractive index and extinction coefficient of the films were measured using spectroscopic ellipsometry (350 − 1150 nm), ellipsometry (633 nm), and reflectance measurements (300 − 1200 nm).

• Electrical Properties: Conductivity observed in some TiO2 films that had undergone a reaction was determined using a four-point probe (FPP). The degree of surface passiva-

tion afforded by TiO2 films and TiO2/SiO2 stacks was determined using the transient photoconductance decay (transient-PCD) technique (discussed further in Chapter 5).

• Chemical Properties: The chemical resistance of TiO2 thin films were determined by placing films in various solutions commonly used in solar cell processing.

A brief introduction to each of these characterisation techniques will be provided here and the results presented throughout this section.

4.2 FTIR Spectroscopy

FTIR spectra provide information regarding the bonding between atoms in a sample, and are primarily used in the semiconductor industry to determine the existence of dopant or impurity atoms. The nature of the technique is very quantitative in identifying the impurity type, but very qualitative in determining the concentration of the impurity. FTIR spectra are also demonstrated in this work to be useful for understanding changes occurring in TiO2 films during high-temperature processing. For a discussion of the theory of FTIR refer to works by Nakamoto257 and Griffiths and de Haseth.258 4.2 FTIR Spectroscopy 93

FTIR measurements in this work were performed with a Nicolet 520 spectrometer in the range 250 − 5000 cm−1.A4cm−1 resolution was used and 256 scans were collected per spectrum. Any oxygen and carbon in the float-zone (FZ) wafers was accounted for by first performing a reference spectrum with only the FZ wafer. This reference spectrum was then subtracted from subsequent sample spectra to yield information about the TiO2 film.

Erkov et al. published some excellent FTIR spectra of 110 nm thick, LPCVD deposited TiO2 thin films, as shown in Figure 4.1(a) and (b).117 The rutile films are deposited on Si wafers ◦ at 630 C using a TiCl4 precursor with N2O and H2 ambient gases. The spectra shown in

Figure 4.1(a) are the as-deposited TiO2 film (curve 1), samples annealed in a vacuum (curve

2), structures annealed in N2O ambient (curve 3), and samples that had 10 nm of SiO2 grown prior to TiO2 deposition (curve 4). In Figure 4.1(b) the IR spectra of a bare Si wafer with 0.4 − 0.6 nm natural oxide (curve 5) and of sample 3 after etching off the TiO2 in a hot −1 H2SO4 etch (curve 6). The strong absorption peak at 608 cm is inherent to the rutile phase, however this is somewhat obscured by a weaker silicon absorption peak at 608−610 cm−1,as seen in curve 5. The overlapping peaks at 423 cm−1 and 460 cm−1 are also characteristic of −1 the rutile modification of TiO2. An additional absorption peak is observed at 470−480 cm for samples annealed in N2O (curve 3) and for structures with a 10 nm interfacial SiO2 layer −1 (curve 4). In the former case the peak at 480 cm is attributed to the formation of Ti2O3. This will be discussed in more detail in Section 5.2.2. The small peak at 515 cm−1 has been observed for TiO2 samples annealed in vacuum only (curve 2). The existence of a peak −1 at 1108 cm is indicative of the formation of SiO2 at the Si:TiO2 interface (curves 2, 3, 4 and 6). This peak is even observed in the vacuum (10−2 Torr) annealed samples where the availability of oxygen would be extremely limited.

Figure 4.2 shows the FTIR spectra of two USD-deposited TiO2 films from this work (TO-11 and TO-12), over the range where the spectra of coated wafers differ markedly from the bare silicon reference wafer. It can be seen that at wavenumbers greater than 570 cm−1 the features observed in all spectra are similar and originate from the silicon substrate. A broad −1 absorption peak at 1083 cm is the only difference between the TiO2 coated wafers and the bare substrate. This peak is likely to be due to the formation of SiO2 at the TiO2:Si interface, and is normally observed at 1108 cm−1.117

When comparing these thin film FTIR spectra to reference spectra for bulk anatase and 259, 260 rutile TiO2 obtained from the literature it is immediately apparent that there is little agreement between the absorption peaks. Only a local transmittance peak at about 380 cm−1 can be observed in all samples. The spectra of the bulk TiO2 samples may differ greatly due to other incorporated impurities or the measurements may not have been obtained at room temperature. TiO2 samples TO-11 and TO-12 were loaded into the furnace under different conditions, and the effect of the ambient furnace gas will be discussed in detail in Section

5.2.2. The spectrum of TiO2 film TO-11 has been successfully modelled by the author using the polarisation-dependent model recently applied to single crystal anatase.261, 262 The model 94 4. Characterisation of TiO2 Thin Films

Figure 4.1: (a) and (b) FTIR spectra of TiO2 thin films on silicon wafers. The spectra are of 1) an as-deposited TiO2 film, 2) TiO2 films annealed in a vacuum, 3) TiO2 films annealed in N2O, 4) TiO2 films deposited on 10 nm of SiO2, 5) a bare Si wafer with 0.4 − 0.6 nm natural oxide, and 117 6) the spectra of sample 3 after etching off the TiO2.

presented in Equation 4.1 is based on the factorised form of the complex dielectric function ε(ν).261 ε2 − ε2 ıγ ε LOn + LOn ε(ν)=ε1(ν) − ıε2(ν)=ε∞ (4.1) ε2 − ε2 + ıγ ε n TOn TOn The longitudinal optical (LO) and transverse optical (TO) phonon oscillator frequency ν (cm−1) and damping γ (cm−1) values are given in Table 4.1. The complex dielectric constant is related to the complex refractive index by ε = n 2 (refer to Section 2.3.6 for more detail). A similar model has been published in the past for rutile.263 The sample structure used in the anatase model was a 300 µm thick polished silicon wafer with a 74 nm thick TiO2

film deposited onto both surfaces. The dielectric constant of the TiO2 film was determined using a Bruggemann effective medium approximation158 of 55.7% Ec-axis and 44.3% E⊥c- axis. The modelling was performed using the WVASE32 software package.62 The best fit 4.3 Raman Spectroscopy 95 obtained with the model is plotted in Figure 4.2 using the anatase values from Table 4.1. It can be seen that the location of the absorption peaks are in excellent agreement with the experimental results. This result is somewhat surprising as the as-deposited films had been subjected to 90 min of annealing at 950◦C, which should be more than sufficient to observe a transformation from anatase to rutile. The only similar result that can be found in the ◦ literature (for undoped TiO2) is for anatase films, deposited at 330 C, remained anatase after annealing at 850◦C.75 In order to confirm that only the anatase phase was present Raman spectroscopy measurements were performed on the films.

Figure 4.2: Comparison of the FTIR spectra of spray deposited TiO2 thin films and the modelled anatase result obtained using Equation 4.1.

4.3 Raman Spectroscopy

Raman scattering occurs when light interacts with the optical phonons in a material. If the photon gives up part of its energy to the lattice, in the form of a lattice vibration or phonon, the photon emerges with a lower energy. It is also possible for a photon to absorb a phonon and emerge at a higher energy, although these events much weaker. The theory of Raman spectroscopy is discussed in detail in the works of Long264 and Nakamoto.257

Since the intensity of Raman scattered light is very weak, the intense monochromatic beam from a laser is required. The collected signal is usually passed through a double monochro- mator and detected by a photodetector. The Renishaw Model 2000 machine located in the School of Materials Science at UNSW is equipped with three lasers, with wavelengths of 514.5 nm, 632.8 nm and 780 nm. The system uses a prism to disperse the signal onto a CCD 96 4. Characterisation of TiO2 Thin Films

Table 4.1: LO and TO phonon frequencies for anatase and rutile to fit di- electric function in Equation 4.1 (adapted from Gonzalez,261 and Gervais and Piriou263).

Anatase Rutile Mode Frequency Damping Frequency Damping ν (cm−1) γ (cm−1) ν (cm−1) γ (cm−1) E c-axis TO 367 68 172 76 (A2u) LO 755 79 796 38 E⊥ c-axis TO 262 36 189 27 (Eu) LO 366 4.1 367 10 TO 435 32 381.5 16.5 LO 876 33 443.5 21.5

ε∞(Ec-axis) = 5.41 ε∞(Ec-axis) = 7.8 ε∞(E⊥c-axis) = 5.8 ε∞(E⊥c-axis) = 6.0

array. While this makes measurements very fast it also limits the accuracy of the instru- ment to about 1.6 cm−1. An additional complication with the Renishaw machine is that the absorption filter designed to remove the laser line at 0 cm−1 begins absorbing at about 200 cm−1. This interferes with the dominant peak of anatase which lies at 143 − 144 cm−1, and rutile also has a smaller peak that lies at the same frequency. Figure 4.3 illustrates the 265 Raman spectra for the three crystalline phases of TiO2 as well as the amorphous phase.

Figure 4.3: Raman spectra of the various phases of TiO2, obtained from powder samples.265 4.3 Raman Spectroscopy 97

It can be seen that the peaks from each TiO2 phase are clearly separated in frequency and therefore easily distinguishable. Table 4.2 presents the Raman peak assignments for single crystal anatase and rutile, and also silicon. Silicon is included as the absorption depth of even the most absorbing TiO2 films deposited in this work is just over 1 µmatλ = 514.5 nm. Since the films are typically 70 nm thick, photons interact with the silicon, which has a similar absorption depth to TiO2 at this wavelength, creating Raman scattering events. This is confirmed in the literature where the Raman peaks of silicon were still observed at 300 cm−1 −1 and 520 cm when performing measurements on 700 nm thick TiO2 films deposited onto 266 silicon wafers using a laser of 532 nm. Thus, it is not possible to observe the A1g and B1g modes of anatase for a thin film deposited onto silicon.

Table 4.2: Raman active phonon frequencies for anatase and rutile (adapted from Ohsaka et al.267 and Gonzalez 261).

Anatase Rutile Mode Frequency Mode Frequency ν (cm−1) ν (cm−1) Eg 144 B1g 143 Eg 197 Eg 447 B1g 399 A1g 612 A1g 514 B2g 826 B1g 514 Eg 639

Figure 4.4 shows the behaviour of the Raman spectra during annealing to temperatures for a fixed time of 40 min.265 At about 250◦C, the 141 cm−1 anatase peak appears and by 425◦C the conversion to anatase is complete.265 The anatase to rutile transformation is clearly observable at 800◦C, and at 1000◦C the conversion to rutile is complete.

Raman measurements from this work are given in Figure 4.5. Note that the data for the 520 cm−1 silicon peak has been omitted for clarity. Samples 1 and 2 were deposited by APCVD at Eurosolare S.p.A., Italy. Sample 1, deposited at 320◦C, exhibits no anatase peaks. Silicon peaks at 300 cm−1 and were observed, and the peaks at 622 cm−1 and 834 cm−1 could possibly be attributed to the A1g and B2g modes of rutile. The A1g peak is typically very strong, while the B2g mode is very weak. Therefore it is concluded that the sample is predominantly amorphous with a very small fraction of rutile. With increased temperature (450◦C) anatase peaks emerge at 143 cm−1,396cm−1 and 637 cm−1. The location of these peaks is in excellent agreement with the bulk values given in Table 4.2. Samples 3 and 4 were both deposited at 450◦C using ultrasonic spray pyrolysis, and sample 4 received a subsequent 90 min anneal at 950◦C. These results confirm the earlier result that only the anatase phase is present after a lengthy high-temperature anneal. The broad peak at about 98 4. Characterisation of TiO2 Thin Films

Figure 4.4: Raman spectra measured after annealing TiO2 powder sam- ples at various temperatures.265

950 cm−1 is attributed to silicon-oxygen-titanium (Si-O-Ti) bond formation.268 Note that a peak at 960 cm−1 is observed in silicon, however this is a multiple phonon event and it is 8 very weak. The Si-O-Ti peak indicates an interaction between the TiO2 layer and either 266 a thin or native SiO2 layer. This result is consistent with the FTIR finding that a SiO2 layer is formed at the TiO2:Si interface. Fitzgibbons et al. noted that TiO2 depositions on quartz substrates remained anatase after annealing at 1000◦C for 20 hr.67 Therefore, it is postulated that the presence of an interfacial SiO2 layer has resulted in the TiO2 maintaining the anatase preferentially.

4.4 X-ray Photoelectron Spectroscopy

X-ray photoelectron spectroscopy (XPS) is primarily used to identify chemical species at the surface of a sample. When high-energy photons (X-rays) interact with atoms of the sample via the photoelectric effect, electrons are ejected from the core levels. However only photoelectrons ejected from atoms in the top 5 − 50 A˚ can escape and therefore be detected. The primary strength of XPS is that it allows chemical and not only elemental identification. Depth profiling is also possible by performing ion-beam sputtering of the sample. Further information on XPS theory and measurement techniques can be found in Briggs and Seah.269

The spectrometers used in this work were, firstly, a VG Scientific ESCALAB 220i-XL, using monochromated Al Kα (1486.6 eV) radiation, with an accuracy of 0.2 at. %. Figure 4.6 shows the XPS spectrum resulting from a wide energy scan of a spray deposited TiO2 film on silicon. This was a surface scan performed with no etching of the film. Apart from the expected titanium and oxygen peaks, a significant carbon peak is observed. This peak is primarily due to adsorbed carbon from the atmosphere, as the sample was stored in 4.5 Rutherford Back-Scattering Spectroscopy 99

Figure 4.5: Raman spectra of TiO2 thin films deposited by APCVD (sam- ples 1 and 2) and ultrasonic spray pyrolysis (samples 3 and 4). Anatase (A) and silicon (Si) Raman peak assignments are shown. Samples 1 and 2 were deposited at 320◦C and 450◦C, respectively, while samples 3 and 4 were both deposited at 450◦C. Sample 4 received a subsequent 90 min anneal at 950◦C. air. This adsorbed carbon had an atomic concentration of 13% at the top surface. As the film was etched this reduced quickly to below 1%, indicating that 1% was the level of contamination resulting from the organic precursor (see Section 5.3.1). The location of the carbon peak at 285 eV serves the useful purpose of providing a calibration point for the spectra, accounting for any charging effects.94 Windows around the titanium, oxygen, carbon and silicon peaks were drawn and examined in greater resolution. The O1s peak at 531 eV exhibits a slight shoulder at exhibited a shoulder at about 533 eV, and this has been attributed to the adsorption of water vapour onto the top surface.94 The results of further XPS experiments will be discussed in Section 5.3.1.

4.5 Rutherford Back-Scattering Spectroscopy

Rutherford back-scattering spectroscopy (RBS) is based on bombarding a sample with high- energy helium ions and measuring the energy of these back-scattered ions. With this method, the masses of elements and their depth distribution in the sample can be determined. The depth resolution of RBS is typically 100 A,˚ and atomic concentrations down to 10 − 100 ppm can be detected. Figures 4.7(a) and (b) indicates how the measured RBS spectrum corre- sponds to the elements in the sample structure.270 Since Au, Ag and N are only on the surface in Figure 4.7(a), the RBS signals have a narrow spectral distribution. It also demon- 100 4. Characterisation of TiO2 Thin Films

Figure 4.6: XPS surface scan over a wide range of binding energies of a TiO2 film spray deposited onto a silicon wafer.

strates that, firstly, the RBS yield increases with atomic number. Secondly, the RBS signal of elements lighter than the substrate (nitrogen in our example) will appear superimposed on the substrate signal, while heavier elements will be displayed as separate peaks.270Figure 4.7(b) demonstrates how the RBS signal for broadens as ions back-scattered from deeper within the gold film have lower energies due to additional energy loss mechanisms within the film. More detailed theory on RBS can be obtained from Schroder270 or Breese et al.271

The RBS results in this work were obtained with a collimated beam of 4He2+ ions at an energy of 2 MeV. The beam was scanned over an area of 1 × 1mm2. Backscattered ions were detected with a surface barrier detector with a solid angle of 28 msr at a scattering angle of 145◦. The beam was provided by the University of Melbourne 5U pelletron accelerator.

Figure 4.8 shows the RBS spectrum obtained from a TiO2 film deposited by ultrasonic spray 272 pyrolysis onto a silicon wafer and a RUMP simulation of a 100 nm thick TiO2 film. As discussed in Section 2.2.5, the density of the film can also be determined from the RBS spectrum. The areal densities of the titanium and oxygen in the film in Figure 4.8 were 17 2 17 2 ρareal =1.51 ± 0.01 × 10 atoms/cm ρareal =3.1 ± 0.2 × 10 atoms/cm , respectively. The resulting stoichiometry is determined to be 2.02 ± 0.13, being limited by the noise in the oxygen statistics.273 Further results from RBS will discussed in Chapter 6. 4.6 Ellipsometry 101

Figure 4.7: (a) Calculated RBS spectrum for a silicon sample with the elements gold (Au), silver (Au) and nitrogen (N) on the top surface. Note the narrow peaks. (b) Au depth profile information corresponds to an increased Au peak width.270

Figure 4.8: RBS spectrum of a TiO2 film on a silicon wafer (solid line) 272 and a RUMP simulation of a 100 nm thick TiO2 film.

4.6 Ellipsometry

4.6.1 Overview

Ellipsometry is primarily used in the semiconductor industry for accurately measuring the thickness of thin dielectric films. However, ellipsometry is also a very powerful tool for determining the optical constants of a film. The following brief introduction to ellipsometry has been adapted from J.A. Woollam.62 102 4. Characterisation of TiO2 Thin Films

Ellipsometry measures the change in the state of polarisation of light that is reflected from the front surface of a sample. The measured ellipsometric parameters Ψ and ∆ are related to the ratio of the Fresnel coefficients Rp and Rs for p- and s-polarised light, respectively. Figure 4.9 shows an incoming linearly polarised beam, with the p-direction lying in the plane of incidence (the plane that contains the incident and reflected beams) and the s-direction (from Senkrecht, German for perpendicular) lies perpendicular to the p-direction.

Figure 4.9: An ellipsometry experiment, showing the p- and s-directions and the electric field E (from J.A. Woollam62).

Figure 4.10 compares linearly, circularly and elliptically polarised light. By looking at the electric field vector E in a plane perpendicular to the direction of propagation the polarisation of the light can be determined. For linearly polarised light (a), E lies in one line at all times and there is no phase difference. In the case of circularly polarised light (b) the Ex and Ey components of E are of equal magnitude but 90◦ out of phase. Circularly polarised light may precess either clockwise or counter-clockwise around the circle. In general, the Ex and

Ey-fields do not have to be of equal magnitude and could possess any phase relationship. In this case (c) the tip of the electric field vector E traces out an ellipse as a function of time. The ratio of the reflection coefficients Rp and Rs is defined as the complex reflection ratio ρ, which is related to the measured ellipsometric parameters Ψ and ∆ by R ρ p ı∆ = = tan(Ψ) exp (4.2) Rs

The schematic diagram in Figure 4.11 shows the typical componentry of a simple ellipsome- ter. A collimated beam of unpolarised light becomes linearly polarised after passing through the polariser (P). The compensator (C), or retarder, changes this linearly polarised light the elliptically polarised light. The compensator contains a fast and slow optical axis per- pendicular to the direction of transmission. Therefore one component will become retarded in phase relative to the other component. After being reflected off the front surface of the sample the linearly polarised light enters the analyser (A) (similar to the polariser). With 4.6 Ellipsometry 103

Figure 4.10: Diagram looking into the propagating beam, showing (a) linearly, (b) circularly and (c) elliptically polarised light (adapted from J.A. Woollam62). null ellipsometry the signal is extinguished by the analyser and a zero output is observed at the detector. Then by fixing certain angles the number of solutions is reduced down to one pair of Ψ and ∆. There are other ellipsometric configurations (e.g., rotating analyser like the VASE instrument) and the reader is referred to J.A. Woollam62 for further information. In the example shown in Figure 4.11 light is reflected at an air-substrate interface. In this

Figure 4.11: Schematic diagram of a typical ellipsometry setup (adapted from270). simple case Fresnel’s equations can be used to show that the complex refractive index n can be determined from the measured Ψ and ∆ values:270

n n − ık = 1 1

4ρ 2 = n0 tan(φ) 1 − sin (φ) . (4.3) (1 + ρ)2 104 4. Characterisation of TiO2 Thin Films

If the ellipsometric ratio ρ from Equation 4.2 is measured at the incident angle φ and n0 is known (unity for air) then the refractive index n and extinction coefficient k can be calculated.

In Figure 4.12(a) and (b) the behaviour of the p- and s-components of the reflectance and the ellipsometric parameters Ψ and ∆ are plotted for a bare crystalline silicon substrate for light of λ = 633 nm. It can be seen that as the angle of incidence increases Ψ reaches a minimum at 75.5◦. At this angle ∆ changes from very close to 180◦ to 0◦. This angle is known as the Brewster angle and is determined by

n −1 1 φB =tan ,(4.4) n0

where φB is the Brewster angle (complex for an absorbing substrate), n0 is the ambient refractive index (air in our example) and n1 is the complex refractive index of the substrate (silicon). The absorption in the silicon substrate prevents the p-polarised component and Ψ from going to zero.

In summary, two important advantages of ellipsometry are that, firstly, because the ratio of two numbers is measured it is highly accurate and reproducible and, secondly, the phase information ∆ makes the measurement very sensitive.

4.6.2 Ellipsometers

Three different ellipsometers were used in this work. A single wavelength G¨artner L116A ellipsometer equipped with a He-Ne laser (632.8 nm) was used for quick determination of film thickness and refractive index immediately after film deposition or high-temperature processing. It was attempted to measure all films at both 50◦ and 70◦ as a consistency check. However, this often was not possible due to the films acting as an AR coating a reducing the signal significantly. Instead, three to five measurements were made at 70◦ and these values averaged.

The second ellipsometer equipped with three lasers (Pacific Solar Pty Ltd, Sydney), with wavelengths of 632.8 nm, 831.7 nm and 1299 nm. This enabled a few points on the dispersion curve of n and k to be measured. Additionally, at each wavelength the sample was measured at incident angles from 45◦ to 80◦ in 5◦ steps providing a high degree of confidence in the extracted thickness and optical constants.

Variable angle spectroscopic ellipsometry (SE) measurements were performed in the De- partment of Physiology at the University of Western Australia with a VASE machine (J.A. Woollam Co., Inc.). Traditionally, a monochromatic light source such as a He-Ne laser is used as a source, however in the case of SE a xenon lamp and a monochromator are imple- mented instead. The ability to measure data a wide range of wavelengths (250 − 1700 nm) 4.6 Ellipsometry 105

Figure 4.12: Behaviour of (a) p- and s-polarisation components, and (b) Ψ and ∆ for a bare silicon substrate (adapted from J.A. Woollam62).

and at multiple angles means that SE is a very powerful technique for determining the op- tical constants of single or multilayer samples. For the TiO2 on silicon samples measured in this work the wavelength range was set to 350 − 1150 nm due to a noisy signal outside this range. The noise in the UV was due to the limited signal from the lamp, while the noise above 1150 nm arose from the type of fibre optic cable implemented. The wavelength range 350 − 1150 nm is excellent for silicon solar cells, with little energy from the sun at less than 350 nm and silicon has an extremely low absorption coefficient above 1150 nm. Data were collected for each sample at angles of 65 − 80◦ in 5◦ steps. 106 4. Characterisation of TiO2 Thin Films

4.6.3 Lorentz Oscillator Model

All modelling of optical constants was performed using the WVASE32 software package.62 The Lorentz oscillator model was chosen for fitting the dispersive optical constants to the ellipsometric data. The Lorentz oscillator model is based on the assumption that the response of electrons in a material to the light beam is similar to the response of a harmonically driven mass on a spring subject to a force (friction).62 In this analogy, the mass represents the electron, the spring corresponds to the electrostatic forces on the electron (due to all the other electrons and nuclei in the material), and the friction represents the electron energy loss due to the emission of a photon. The imaginary part of the dielectric function ε2 is proportional to the power per unit volume absorbed from a monochromatic light source of a given wavelength.62 To apply this to the mass on the spring analogy for the electron, the power absorbed by the mass from the driving force is then calculated using electromagnetic units. After determining ε2 the Kramers-Kronig transformation can be applied to obtain the real part of the dielectric function ε1. The relationships between the dielectric constants

ε1 and ε2 and the refractive index n and extinction coefficient k are

2 2 ε1 = n − k (4.5)

ε2 = −2nk (4.6)

The Lorentz oscillator model is Kramers-Kronig consistent, and is usually expressed in terms of the dielectric constant, where

ε(E)=ε1 − ıε2 N Ai Bi Eni = ε1(∞)+ . (4.7) E2 − E2 − ıB E i=1 ni n

In Equation 4.7, the incident photon energy E has the units of eV, ε1(∞) is the value of the real part of the dielectric function at very large photon energies (dimensionless), Ai is th th the amplitude of the i oscillator (dimensionless), Bi is the broadening of the i oscillator th (eV), and Eni is the centre energy of the i oscillator (eV). Initial modelling was performed with a single Lorentz oscillator, however it was subsequently found that a double Lorentz oscillator was able to model the experimental behaviour of some TiO2 films much better. These situations included when the measurements approached the bandgap of anatase or rutile, or when there was an anatase/rutile phase mixture present in the sample. Other researchers have made use of similar single and double oscillator models to describe the dielectric function of a material.45, 140, 144, 151

WVASE32 uses the mean-squared error (MSE), based on the chi-square (χ2) test, to deter- mine the quality of match between the data calculated from current model (ψmod,∆mod)and 4.6 Ellipsometry 107 the experimental data (ψexp,∆exp). The MSE is defined in Equation 4.862   2 2 1 N ψmod − ψexp ∆mod − ∆exp MSE =  i i + i i  (4.8) 2N − M σexp σexp i=1 ψ,i ∆,i 1 = χ2 , 2N − M where N is the number of (ψ,∆)pairs,M is the number of variables in the model, and σ are the standard deviations on the experimental data points. The MSE should reach an absolute minimum for the best fit, and this minimum should be fairly sharp. However, insensitive parameters and correlation between parameters can hinder the determination of an absolute minimum MSE. It is therefore necessary to spend a significant amount of time developing the initial model for the deposited material and substrate. Once established, it is then easier to apply this model to similar films with greater confidence.

4.6.4 Surface Roughness Model

The most common method of modelling surface roughness of TiO2 thin films is to use a Bruggeman effective medium approximation (EMA), comprised of 50% void and 50% 82, 128, 161, 274 TiO2. The expression for the Bruggeman EMA was given in Equation 2.10 in

Section 2.3.6. The TiO2 fraction of the EMA was coupled to the underlying dense TiO2 layer, linking the optical constants of the surface roughness layer to the denser TiO2 layer below. The only parameter permitted to vary in the EMA layer was its thickness. This approxi- mation for surface roughness has been used successfully by other researchers for modelling 82, 128, 275 82, 275 TiO2 films. Mardare and co-workers noted that surface roughness significantly changed the values of the optical constants, and suggested independently confirming the sur- face layer roughness using atomic force microscopy (refer Section 4.9). Furthermore, it was stated that for large surface roughness a graded layer model becomes necessary. The author evaluated this option however it was found that the increased number of fitting parameters easily lead to unphysical results. The results using either a single or double Lorentz oscillator coupled with the 50% void/50% TiO2 EMA layer were found to be entirely satisfactory.

4.6.5 Ellipsometric measurements of Spray Deposited TiO2 Thin Films

◦ The optical constants of a TiO2 film spray deposited at 450 C are plotted in Figure 4.13. A ◦ ◦ second TiO2 film deposited at 450 C but annealed at 950 C was also measured. The optical constants were determined using the three-wavelength ellipsometer. As the optical constants in the region 632.8 − 1299 nm are fairly non-dispersive the reflectance of the same samples was also measured using a ultraviolet-visible-near infrared (UV-vis-NIR) spectrophotometer 108 4. Characterisation of TiO2 Thin Films

(see Section 4.7). A Lorentz single-oscillator model was used for both ellipsometric and reflectance data. It can be seen from Figure 4.13 that the agreement in the refractive indices is very good, while there is some discrepancy in the extinction coefficient data (note that k is on a log scale). The surface roughness of these films is large (refer to Section 4.9) and, as discussed previously in Section 2.3.1, surface roughness can affect measurements of the extinction coefficient significantly. Additionally, it is known that ellipsometry is very sensitive to non-uniformities in both thickness and refractive index.276 For these reasons, the reflectance data is believed to provide the best estimate of the behaviour of the extinction coefficient of the spray deposited TiO2 films.

Figure 4.13: Comparison between the optical constants of as-deposited and annealed TiO2 films using three-wavelength ellipsometry (symbols) and reflectance (curves) measurements. The hollow symbols (2,◦)are refractive index values while the solid symbols (,•) represent extinction coefficient values.

The refractive indices of the spray deposited films before and after annealing (at λ = 600 nm) are 2.449 and 2.458, respectively. The high refractive indices observed here are possibly linked to the high deposition rate of the spray system.277 Additionally, the high refractive index of the as-deposited film indicates that they are very dense. Equation 2.2 from Section 2.2.5 can be used to estimate the densities of the two films - the refractive indices of 2.470 and 2.476 (at λ = 550 nm) correlate to densities of 3.63 g/cm3 and 3.64 g/cm3, respectively. This compares to a density of 3.84 g/cm3 for single crystal anatase. These results can be compared to other

TiO2 film densities reported in the literature. Ottermann et al. deposited stoichiometric anatase films using e-beam evaporation at 320◦C.278 These films had a refractive index of 2.345 at 550 nm and a density of 3.3 g/cm3. Bendavid et al. measured the density of filtered-arc deposited anatase thin films to be 3.82 g/cm3 with a refractive index of 2.62 at 4.7 Reflectance Spectrophotometry 109

550 nm.33 This result is interesting in that from Equation 2.2 a refractive index of 2.62 should correspond to a density of 3.98 g/cm3, which is greater than the bulk anatase value. Obtaining a density greater than the bulk phase would require Ti and O atoms to occur interstitially within the anatase crystal. This would most likely result in a fraction of the film converting to the more stable phase of rutile. In the X-ray diffraction plot indicating the anatase phase a small peak for rutile(110) is visible. This small rutile fraction could explain the high film density.

Furthermore, based on the results from Sections 4.2 and 4.3, which indicated that the films are 100% anatase, the porosity of the films can then also be determined using Equation 2.3, with the bulk anatase value nb being 2.532 at 600 nm. This calculation results in a film porosity of 7.6% for the as-deposited sample and 6.8% for the annealed sample.

4.6.6 SE measurements of CVD TiO2 Thin Films

The thin films deposited via CVD had much smoother surfaces (refer to Section 4.9). Addi- tionally, the restriction on the range of deposition temperatures was relaxed, and depositions ◦ could be performed anywhere from 150−450 C. This enabled the refractive index of the TiO2 films to be varied. The presence of water vapour during depositions was also investigated. The results of these experiments are presented and discussed in detail in Chapter 7.

4.7 Reflectance Spectrophotometry

Reflectance spectrophotometry is typically employed for measuring the thicknesses and op- tical constants of deposited thin films. The theory of the method is well developed and ex- 38, 279, 280 cellent explanations can be found in the literature. The reflectance spectra of TiO2 films were measured using a Varian Cary 5G ultraviolet-visible-near infrared (UV-vis-NIR) spectrophotometer equipped with a Labsphere DRA-50 integrating sphere. This enabled re- flectance measurements in the wavelength range 250−2500 nm to be performed. Reflectance measurements were performed for determining the optical properties of spray deposited TiO2 thin films (see Section 4.6.5) where ellipsometric determination at short wavelengths was not possible. Additionally, the optical constants of multilayer TiO2 films (see Chapter 7) were determined using both spectroscopic ellipsometry and reflectance spectrophotometry.

4.8 Electron Microscopy

Field emission scanning electron microscopy (FESEM) digital images of TiO2 films were recorded using a Hitachi S-900 instrument with an accelerating voltage of 10 − 12 kV. Even 110 4. Characterisation of TiO2 Thin Films though the films were insulating, it was not necessary to coat the films with a thin layer of chromium. Images of USD-, APCVD- and CVD-deposited TiO2 thin films will be shown in the following sections.

4.8.1 USD-Deposited TiO2 Thin Films

Cross-sectional SEM images of USD-deposited TiO2 films on silicon substrates are shown in

Figure 4.14(a), (b) and (c). Figure 4.14(a) shows an as-deposited TiO2 film. Figure 4.14(b) indicates the possible existence of voids underneath the TiO2 film. The uniform, thin film on top of the silicon substrate is SiO2. The formation of this film will be discussed in Chapter 5. The image in Figure 4.14(c) is the same film as in (b), however the magnification is three times greater at 600,000 times. In all the Figures it can be seen that the USD-deposited films appear to be continuous and dense. This is in agreement with the high refractive indices ◦ measured in Section 4.6.5. Figure 4.15 shows a plan-view of an as-deposited (450 C) TiO2 film. Again, the crystals appear to be densely packed, and the possibility of some sintering having occurred is not ruled out.

4.8.2 APCVD-Deposited TiO2 Thin Films

The SEM image in Figure 4.16 indicates the presence of particulates, both incorporated into andontopoftheTiO2 film, as well as cracks in the film. Cracks were only observed in films deposited using a commercial APCVD system, and the limited and infrequent access to the deposition system meant that it was not possible to determine their origin. Particulates were observed in both USD- and APCVD-deposited films. The particulates shown in Figure 4.16 are representative of smaller particulates, with larger sizes (up to 30µm) being observed in USD-deposited samples. The inability to deposit TiO2 film without the presence of particulates was a major motivating factor for developing the CVD-based system.

4.8.3 CVD-Deposited TiO2 Thin Films

Figure 4.17 contains four plan-view SEM images of TiO2 thin films, all deposited via CVD at 450◦C. In (a), angular crystallites can be seen, many of which are separated by large voids. It is immediately apparent that the films deposited via CVD are less dense than the USD-deposited ones. The lower-density of the as-deposited CVD films was reflected in the refractive index measurements in Chapter 7. After annealing for one hour (b), the agglomeration of grains are observed and the edges of the crystallites have become rounded. Annealing at a slightly higher temperature for a period of 6 hr resulted in the majority of grains sintering together. The largest voids in between the grains remain “unbridged”. Annealing for a further 16 hr at 1000◦C results in an almost continuous film, with small voids 4.8 Electron Microscopy 111

Figure 4.14: USD-deposited TiO2 films on Si substrates: (a) one sample as-deposited, (b) a second sample, indicating the possible existence of voids underneath sections of the film, and (c) a higher magnification image of (b) showing the dense and continuous nature of the film.

(typically 40 nm in diameter) remaining. Figure 4.17(a)−(d) demonstrate that a sintering process has occurred, similar to that illustrated in Section 2.2.1. The densification of the

TiO2 layer is accompanied by a reduction in film thickness, from 79 nm in (a) to 64 nm after step (d). The correlation between density and refractive index is discussed in Section 2.2.5.

Due to the horizontal angle of the impinging vapour it is not possible to coat textured wafers ◦ with this TiO2 deposition method. The result of performing a 1 hr deposition at 450 Cis dramatically shown in Figure 4.18. Large crystal structures over 1 µminsizehavegrownon the tips of the pyramids, while the growth of smaller structures also appeared on the upper regions of the pyramid edges. The structures observed here are very similar to the “grasses” 101 and “plants” observed by Goossens et al., obtained using a TiCl4/TPT precursor mixture. It was noted that these fractal structures are good for solar harvesting in Gr¨atzel-type solar cells, due to their large surface area. It was not attempted to further adjust the system due to the requirement of the horizontally incident vapour in order to maintain the grooves free 112 4. Characterisation of TiO2 Thin Films

Figure 4.15: As-deposited TiO2 film USD-deposited onto a silicon sub- strates. The crystals seem to be relatively densely packed. The darker regions are most likely crystals of a different orientation.

Figure 4.16: SEM image of an APCVD-deposited TiO2 film exhibiting particulates and cracks.

from TiO2.

It was anticipated that the horizontally impinging stream of TPT vapour would result in

TiO2 film deposition on the top surface of a wafer but not in the grooves. To test this 4.8 Electron Microscopy 113

◦ Figure 4.17: SEM images of TiO2 films deposited by CVD at 450 C: (a) as-deposited; (b) 1 hr anneal (load 800◦C and ramp to 950◦C); (c) 6 hr anneal at 1000◦C; and (d) 22 hr anneal at 1000◦C.

hypothesis the grooves and busbar on the wafer were oriented at 45◦ to the incoming vapour, as illustrated in Figure 4.19(a). A new busbar design was implemented in order to keep the groovesfreeofTiO2. This involved increasing the step the laser took in between scribing adjacent busbar grooves, as shown in Figure 4.19(b) and (c).

Figure 4.20(a), (b) and (c) are SEM images of a silicon wafer after performing CVD depo- sition at an incident angle of 45◦. The V-shape groove in Figure 4.20(a) arises from the 114 4. Characterisation of TiO2 Thin Films

Figure 4.18: SEM images of CVD TiO2 films deposited with a horizon- tally impinging TPT vapour onto textured silicon wafers at 450◦C.

E

F

D 737

Figure 4.19: Diagram depicting: (a) the 45◦ orientation of the busbar and fingers to the impinging TPT vapour; (b) the cross-section of the new busbar, designed to keep the grooves free of TiO2; and (c) a traditional BC solar cell busbar.

NaOH based anisotropic groove etch used. Large overhanging TiO2 deposits can be seen on the top-edges of both sides of the groove in Figure 4.20(b) and (c). These deposits are up to 3 µm in length. It is believed that this will not inhibit electroless metal plating and, in fact, a previous work may indicate that the overhanging dielectric may help promote plating at the depths of the groove rather than near the surface.281

An energy dispersive X-ray spectrometry (EDS) attachment to a SEM enables detection and 4.9 Atomic Force Microscopy 115 identification of X-rays produced by the impact of an electron beam on the sample. The detection limit of EDS is approximately 0.1 − 1.0 at.%. The EDS spectra of a TiO2-coated BC solar cell are given in Figure 4.21. The spectra from the top surface exhibits a large peak at 1.75 keV (Si), a small Si sum-peak at 3.5 keV, and a large Ti peak at 7.5 keV. The Si sum-peak results from two Si atoms hitting the detector simultaneously, and therefore an energy of 3.5 keV is measured. The Si sum-peak is typically 1000 times smaller than the main Si peak at 1.75 keV. In the groove spectrum, the height of the Ti peak is reduced to roughly four times the height of the Si sum-peak. Although EDS is essentially a qualitative measurement, the relative heights of the Si sum-peak and the Ti peak can provide some indication as to the abundance of Ti in the grooves. Therefore, some TiO2 was observed in the grooves, however it is believed that this amounts to a fraction of a percent. The peaks at 4.5keVand4.95 keV belong to nickel (Ni). Thus, as well as plating in the heavily-doped grooves, Ni is also plating either on top of the TiO2 or is contacting the doped emitter through the TiO2.

Figure 4.20: (a), (b) and (c): SEM images of large TiO2 deposits at the edges of laser-scribed grooves.

4.9 Atomic Force Microscopy

Atomic force microscopy (AFM) was performed in order to determine the surface rough- ness of the TiO2 films. The surface roughness can be quantified by the root-mean-squared 116 4. Characterisation of TiO2 Thin Films

Figure 4.21: EDS spectra from the top surface and the grooves of a TiO2 coated BC solar cell after electroless nickel plating.

roughness (RRMS), which is defined as the standard deviation of the AFM data,

N (z − z¯)2 R = n=1 n (4.9) RMS N − 1

th where zn is the height of the n data,z ¯ is the mean height and N is the number of data. The measurements were performed with a Digital Instruments Nanoscope III microscope (Electron Microscope Unit, UNSW) in contact mode.

Figure 4.22 depicts the TiO2 surface morphology of an as-deposited CVD film (a), an an- nealed CVD film (b), and a USD-deposited film (c) over a 1.5 µm × 1.5 µm area. The sample in (a), deposited at 450◦C, is relatively smooth for a polycrystalline film and has an surface roughness of RRMS =4.9 nm. The surface roughness is observed to increase upon annealing ◦ (1 hr at 950 C) to RRMS =7.9 nm due to growth in the grain sizes. The samples in Fig- ure 4.22 (a) and (b) are the same samples depicted in Figure 4.17(a) and (b). The flatter regions of the annealed sample observed with SEM (see Figure 4.17(b)) are not visible in ◦ the AFM image. Figure 4.22 (c) depicts a TiO2 film USD-deposited at 450 C. The surface roughness for this film is RRMS =10.6 nm, and a large variation in film thickness can be seen. Over a larger area (85 µm × 85 µm) the surface roughness increased dramatically to

RRMS =32.9 nm and the occasional deep pinhole could be seen. 4.10 Chemical Resistance 117

Figure 4.22: AFM images of (a) as-deposited CVD TiO2 films, (b) CVD- ◦ ◦ deposited TiO2 film after 1 hr anneal (load 800 C and ramp to 950 C), and (c) a USD-deposited TiO2 film.

4.10 Chemical Resistance

Both USD- and CVD-deposited TiO2 films were placed in various chemical solutions. The chemical resistance of the films were ascertained by observing any colour change and etching that occurred before and after etching. Table 4.3 summarises the results for both USD- and

CVD-deposited TiO2 films and the time spent in solution. In Table 4.3, a tick (✓) indicates that no change in colour or the appearance of pinholes could be detected, while a cross (✗) means that the TiO2 films were significantly etched. Additional explanations are included in footnotes below Table 4.3. The exact composition of the solutions can be found in Section ◦ 2.5. CVD films deposited at 450 C were not as chemically resistant as TiO2 films deposited by USD at the same temperature. All non-annealed TiO2 films were susceptible to etching from highly alkaline solutions. This included chemicals like NaOH, NH4OH, and the Cu plating solution (Enplate 704, Melbourne, pH=11). The resistance against acidic solutions was very good except for the HF/HNO3 mixture used in the CP etch, which removed films and began etching the silicon substrate after a few seconds. The green Nickelex solution (Transene Inc., MA, USA) used had a pH of 4. 118 4. Characterisation of TiO2 Thin Films

Table 4.3: Chemical resistance of both as-deposited annealed TiO2 films. Chemical Time in USD CVD CVD Solution Solution (450◦C) (anneal) Saw-damage etch 20 min ✗✗ ✗ Groove etch 20 min ✓ ✗✓ CP etch 2 min ✗✗ ✗ Dilute HF 1 min ✓✓ ✓ BHF 1 min ✓✓† ✓† RCA1 5 min ✗‡ ✗‡ ✗‡ RCA2 5 min ✓✓ ✓ H2SO4 clean 5 min ✓✓ ✓ Nickelex 2 min ✓✓ ✓ Copper > 3hr ✗∗ ✗✓ Some etching of the TiO2 film occurred † Rapid etching of TiO2 film occurred when wet samples were placed in BHF solution ‡ Slight etching of all TiO2 films occurred during RCA1 cleaning, therefore an H2SO4 based cleaning step was used instead ∗ Chemical resistance to basic copper plating solution only achieved after annealing

4.11 Conclusions

Extensive characterisation of the USD- and CVD-deposited TiO2 thin films provided very useful information about the material properties and the behaviour of these properties under high temperature processing. A surprising result was that the Raman spectra of both as- ◦ deposited and annealed (950 C for 90 min) samples indicated that the USD-deposited TiO2 films were anatase. This is the highest temperature reported where anatase films have with- stood the transformation to rutile. The exact reason for the retardation of the anatase-rutile phase transformation is not understood, however it is postulated that the thin interfacial

SiO2 layer, grown at the start of the annealing step, could be inhibiting the phase transfor- mation. The IR spectra of TiO2 thin films were able to be modelled well using a factorised form of the dielectric function. As well as confirming the phase of the USD-deposited TiO2 thin films as that of anatase, the FTIR spectra were able to detect non-stoichiometry in TiO2 films annealed in nitrogen (discussed in detail in Section 5.2.2). Both XPS and RBS spectra rendered excellent results, which will be discussed fully in Chapters 5 and 6, respectively.

Spectroscopic ellipsometry was found to be a very accurate method for determining the optical constants of the TiO2 films. The optical properties could be well described by a single or double Lorentz oscillator model, implemented in the software package WVASE32.62 It was necessary to use a 50% void/50% TiO2 EMA as a surface roughness model. The refractive ◦ ◦ indices of as-deposited (450 C) and annealed (950 C) USD-deposited TiO2 thin films were 2.449 and 2.458 at 600 nm, respectively. The high refractive indices are indicative of a 4.11 Conclusions 119 dense film (3.64 g/cm3). The extinction coefficient of the films remained low at 0.1 or less for wavelengths greater than 350 nm. The optical properties of CVD-deposited films are presented in Chapter 7.

SEM images confirmed the dense appearance of the USD-deposited films, although it is possible that voids exist underneath sections of the film. TiO2 films deposited in an indus- trial APCVD system exhibited cracks and large numbers of particulates. With the simple

CVD system, as-deposited TiO2 films had grains of about 30 nm diameter and the density was noticeably less than USD-deposited films. The density was observed to increase with annealing time, due to a sintering process occurring in the furnace. SEM images and EDS analysis have shown that the horizontally impinging vapour has prevented TiO2 deposition in the grooves. Large overhanging TiO2 depositions were observed on both top-edges of the groove. The surface roughness of USD-deposited films was large at 32.9 nm, however CVD- deposited films were much smoother at 4.9 nm and 7.9 nm, as-deposited and after annealing, respectively.

The chemical resistance of both USD- and CVD-deposited films was excellent against all acids, except the aggressive CP etch. The chemical resistance of the films versus basic solutions was not as good, with slight etching occurring during RCA1 cleaning. The etch resistance of the films to alkaline solutions improved somewhat after annealing. 120 4. Characterisation of TiO2 Thin Films Chapter 5

Enhancing the Passivation of TiO2-coated Wafers

Novel applications for titanium dioxide (TiO2) thin films have the potential to reduce produc- tion costs of high-efficiency commercial silicon solar cells, especially for structures like the buried-contact solar cell. This chapter demonstrates, firstly, that a TiO2 film deposited via a cheap and industrially compatible process and does not contaminate the silicon wafer or fur- nace after lengthy high-temperature processing. It has been determined that spray-deposited

TiO2 films are, however, sensitive to the furnace ambient. Thermochemistry and FTIR anal- ysis confirmed that reduction of the TiO2 film occurred when samples were loaded into a pure

N2 environment and titanium sesquioxide (Ti2O3) was formed.

It is demonstrated that good surface passivation of lightly diffused n-type solar cell emit- ters using TiO2 thin films can be achieved when treated with a furnace oxidation process. Transient-PCD, XPS and SEM measurements indicate that the silicon dioxide layer formed at the TiO2:Si interface provides excellent surface passivation. Emitter dark saturation cur- rent densities of 4.7 × 10−14 A/cm2 are achieved by this method, demonstrating for the first time that TiO2 films are compatible with high-efficiency solar cell structures.

5.1 Introduction

Surface passivation is an extremely important design consideration for high-efficiency solar cells, especially at the front surface where the majority of the light is absorbed. Com- mon methods for Si surface passivation include thermal oxidation at temperatures of about ◦ 1000 C to grow silicon dioxide (SiO2) and plasma-enhanced chemical vapour deposition

(PECVD) of hydrogenated amorphous silicon nitride (a-SiN:H). Although TiO2 thin films are the prevalent antireflection (AR) coating in the PV industry, one shortcoming of is that these films afford very little surface passivation to silicon surfaces. Therefore, one aim of

121 122 5. Enhancing the Passivation of TiO2-coated Wafers this work was to develop a method for enhancing the level of surface passivation achievable using commercially-viable TiO2 deposition techniques. The standard buried-contact (BC) solar cell uses a thermally grown SiO2 layer as its primary dielectric film. The SiO2 film can easily withstand subsequent high-temperature processes, such as diffusion, oxidation and aluminium alloying at temperatures up to and beyond 1100◦C. Therefore, several experi- ments were performed to investigate the suitability of TiO2 thin films to high-temperature processing. These experiments examined the stability of the TiO2 films, and whether high- temperature processing of TiO2 thin films deposited using the 97% pure tetraisopropyl ti- tanate (TPT) precursor would result in contamination of the silicon wafer or furnace.

In this work, transient photoconductance decay (transient-PCD) is used extensively in order to determine the effect that TiO2 thin films, deposited onto the surfaces of silicon wafers, have on the surface and bulk recombination properties of the sample. This technique has become a very popular method for determining the bulk minority carrier lifetime (τbulk)and 282 emitter dark saturation current density (J0e) since its conception in 1985. In the transient- PCD technique, minority carriers are injected into the sample using a flash lamp, and the conductivity of the sample is monitored using an inductively coupling bridge. Thus, the method does not require any contacts to the sample. For measurements made with the bulk region in high injection, the emitter recombination has a quadratic dependence on the excess carrier density, while recombination in the base exhibits a linear dependence. This enables the emitter dark saturation current density J0e to be separated from the bulk recombination or at the surfaces. Thus, for a symmetrical n-type sample with the same n-type diffusion and dielectric on both surfaces and assuming the minority carrier diffusion length is much greater than the wafer thickness, the effective minority carrier lifetime τeff extracted from transient-PCD measurements is related to J0e and τbulk by

1 1 2 J0 ∆p 1 = + e + (5.1) τ τ qn2 W τ eff bulk i Auger where W is the wafer thickness, ∆p is excess minority carrier (hole) concentration, ni is the intrinsic carrier concentration and q the electronic charge. After the Auger recombination term has been subtracted, the J0e can be calculated from the constant slope regions of a plot 282 of 1/τeff versus ∆p.

Lightly-doped (1000 Ω cm) n-andp-type float zone (FZ) wafers were used in this work for J0e and τbulk measurements to minimise bulk recombination. A lightly-doped emitter was used on both surfaces, firstly, as the J0e will be extremely sensitive to any change in recombination at the surfaces,22 which makes it a good indicator as to the level of surface passivation achieved. Secondly, similar lightly diffused emitters are commonly used in the fabrication of high-efficiency solar cells,23, 283 will be the most responsive to enhanced surface passivation schemes. Surface recombination was reduced using either thermally grown SiO2 or P2O5:SiO2 layers for passivation. All experiments were performed in high injection (∆p ≈ 4×1015 cm−3), where the dopant concentration for the n-type and p-typewaferswere4.4 × 1012 cm−3 and 5.2 Stability of TiO2 at High-Temperatures 123

1.3 × 1013 cm−3, respectively.

5.2 Stability of TiO2 at High-Temperatures

5.2.1 Titanium Contamination of Silicon

As silicon solar cells are minority carrier devices, maintaining a high effective minority carrier lifetime is extremely important in order to achieve high conversion efficiency. As can be seen from Equation 5.1, this can be achieved by minimizing recombination at the surfaces, repre- sented by the emitter dark saturation current density J0e, and maximising the bulk minority carrier lifetime τbulk. Ensuring that there are no metal ions present during high-temperature processing steps, for example, is vital to maintain a high τbulk. Even low concentrations of metal ions can diffuse into the silicon wafer and create regions of high recombination that can drastically reduce τbulk.

Therefore, there was initially some concern as to whether the TPT precursor or the TiO2

film would introduce contamination into the furnaces and reduce τeff of the silicon wafers. Of all the 3d transition metals, titanium exhibits the lowest known solubility and diffusivity in silicon,284 however its diffusivity is still several orders of magnitude greater than that of dopant atoms like phosphorus and boron. Due to its strong affinity for oxygen and nitrogen, the diffusion of titanium is hindered by even small amounts of residual air in the furnace ambient.284

Studies have been published measuring the performance degradation of solar cells that were fabricated from silicon ingots that had controlled additions of titanium.285–287 An onset of performance reduction was measured with titanium concentrations of 3 × 1011 atoms/cm3, and at a concentration of 2×1014 atoms/cm3, the efficiency of a p-type solar cell was reduced to 37% of its uncontaminated value.287 Additionally, more than an order of magnitude reduction in the bulk minority carrier lifetime resulted. However, the performance of an n-type solar cell was only reduced to 80% of its initial value at a Ti concentration of 2 × 1014 atoms/cm3 (4 parts-per-billion).285 The solid solubility of titanium in silicon is about 1014 atoms/cm3 at 1200◦Cand1012 atoms/cm3 at 950◦C.284 In the worst case scenario, the performance reduction can be calculated by noting that the solid solubility of Ti in Si is 1.5 × 1012 atoms/cm3 at a typical maximum processing temperature used in a solar cell fabrication sequence (1000◦C). This would result in a reduction in efficiency of about 5% for p-type substrates, but no observable change in efficiency for n-type solar cells. The greater impact on p-type material can be attributed to the high majority-carrier capture cross- sections of titanium donors and acceptors, which have a marked influence on the carrier lifetime in p-type silicon.287 One paper has also been published, indicating that titanium 124 5. Enhancing the Passivation of TiO2-coated Wafers may act as an n-type dopant atom in silicon 1.

Importantly, it is known that TiO2 layers deposited onto SiO2 layers are very stable. Keddie et al. determined that there was no intermixing of TiO2 and SiO2 layers after 10 hr at ◦ 233 450 C, despite significant densification of both films and crystallisation of the TiO2 film.

The interface width was measured to be 0.8 nm. Guenther fabricated multi-layer SiO2/TiO2 coatings using the ion-plating deposition technique and observed that the interfaces were 72, 231 sharp when depositing SiO2 on TiO2, but not in the other direction. It is assumed that the titanium ions possess enough energy during deposition to reduce the previously deposited

SiO2 layer. These diminishing interfaces resulted in a lower than expected transmittance from the multilayer stacks. This phenomenon has not been observed with other deposition methods. Additionally, it has been noted that crystalline TiO2 (either anatase or rutile) and SiO2 do not form a mixed compound, although there is some solubility between the 288 amorphous phases of TiO2 and SiO2.

Experiment

Two MiniBrute quartz tube furnaces were available for annealing the TiO2 films, while a third, larger furnace was equipped with a phosphorus trioxychloride (POCl3) liquid dopant source (Schumacher, CA, U.S.A.). A range of quartz boats enabled high-temperature pro- cessing of 4” and 2” diameter wafers as well as quarters of 4” round wafers and 5 × 5cm2 multicrystalline silicon (mc-Si) wafers. All furnaces were equipped with nitrogen, oxygen, forming gas (4% hydrogen in argon) and trichloroethane (TCA) for cleaning. The furnaces were profiled to have a flat-zone of at least 25 cm long with a temperature variation of ±0.5◦C across the flat-zone.

The diagram in Figure 5.1 depicts the processing steps included for the “contamination”, “reduction” and “surface passivation” experiments performed by the author. The substrates for all experiments were either lightly-doped (1000 Ω cm) n-orp-type float zone (FZ), (100)- oriented wafers to minimise the amount of recombination in the base. For ellipsometry measurements, polished 0.1 − 10 Ω cm n-type FZ wafers were used.

For the ”contamination” experiment in this section, several batches of wafers received a heavier 10 Ω/2 POCl3 diffusion. The samples were heavily diffused in order to be less sensitive to the poor level of surface passivation afforded to the silicon wafer by the TiO2 ◦ layer. These wafers were loaded into the furnace 800 CinO2 before ramping the furnace ◦ temperature up to 950 CinanN2 ambient. The total time spent in the furnace was 2 hr.

1Brown and Grannemann reported that titanium diffused to a depth of 150 nm during a 15 min process ◦ 31 at 1000 C from a reduced titanium dioxide (TiO2−x)film. 5.2 Stability of TiO2 at High-Temperatures 125

“Contamination” “Reduction” “Surface Passivation”

NaOH etch (30%, 85°C, 20 min)

RCA2, RCA1, RCA2, HF dip, DI rinse

POCl3 Emitter POCl3 Emitter Diffusion, 10 Ω/V Diffusion, 175 Ω/V

Dilute HF dip to remove PSG

TiO2 Deposition (both sides, ~70 nm thick, Tdep=450°C)

H2SO4:H2O2:DI (1:1:5), RCA2, DI rinse

Furnace Step: Furnace Step: Furnace Step: oxidation (O2) and oxidation (O2) and anneal (N2) anneal (N2) anneal (N2)

Figure 5.1: Diagram showing the processes steps used for the “contami- nation”, “reduction” and “surface passivation” experiments described in this chapter.

Results and Discussion

Transient-PCD measurements were performed in order to determine whether any degrada- tion in bulk minority carrier lifetime τbulk had occurred due to titanium contamination. It can be seen from the results in Table 5.1 that τbulk remains above 2 ms in all cases. The

J0e is relatively insensitive to the change in surface conditions due to the heavy phosphorus diffusion, and varies from 5.6 − 6.9 × 10−13 A/cm2. FZ wafers with only an emitter diffusion and no TiO2 coating placed in the furnace adjacent to the TiO2 coated samples also exhib- ited similarly high τbulk. This indicates that TiO2 thin films, deposited using the same TPT precursor as used in the PV industry, are compatible with high-temperature processing and do result in contamination of the silicon wafer.

5.2.2 Reduction of TiO2

While performing the contamination experiments, it was noted that the gas ambient that the TiO2-coated wafer was loaded into was critical. This was investigated further to fully 126 5. Enhancing the Passivation of TiO2-coated Wafers

Table 5.1: Transient-PCD analysis demonstrating that no bulk contam- ◦ ination has resulted from the TiO2 film after processing at 950 C for 2 hours.

Transient-PCD τbulk J0e performed after: (ms) (A/cm2) −13 10 Ω/2 POCl3 diffusion 2.4 6.2 × 10 −13 TiO2 deposition 2.2 6.9 × 10 950◦C anneal 2.1 5.6 × 10−13

understand the reactions that were taking place. No changes in the appearance of the TiO2

film were observed if oxygen (O2) was present in the furnace when loading the samples (at ◦ 800 C). However, if the wafers were loaded in pure N2, a reaction between the TiO2 film and the silicon substrate occurred. This could be visually observed as regions with a purplish-blue colour. It was postulated that in the absence of oxygen, the silicon reduces the TiO2 to form a non-stoichiometric sub-oxide TiOx. Other researchers have noted that non-stoichiometric 29, 73, 132–136 TiyOx films have a blue, purple, grey-blue or blue-black colour, depending on the film stoichiometry. Shannon and Pask demonstrated that the grey colour appeared in rutile

(the high temperature crystalline phase of TiO2) films once the O:Ti ratio (x)decreased from 2 to 1.991.88 Due to the increased number of oxygen vacancies, the optical absorption 43, 58 in the visible is dramatically increased for TiOx films. It is imperative to minimise the amount of absorption originating from the film for it to act as an efficient antireflection (AR) coating for a solar cell.

Thermochemistry Analysis Reactions that take place between thin films and the substrate may be better understood by considering the thermochemistry of the deposited film in its ambient condition. The standard free energy change of a reaction ∆G◦ is given by G◦ G◦ − G◦ ∆ = ∆f Products ∆f Reactants ,(5.2)

◦ where ∆f G is the standard free energy of formation for 1 mol of a compound from its constituent elements in standard states and is given by the Gibbs-Helmholtz equation (at a constant temperature)

◦ ◦ ◦ ∆f G =∆f H − T ∆S .(5.3)

◦ ◦ −1 −1 In Equation 5.3, ∆f H and ∆S are the change in enthalpy (kJ mol ) and entropy (J mol −1 ◦ K ), respectively, from O K to the reaction temperature T (K). Once ∆f G is calculated for each product and reactant in the system, the ∆G◦ for the whole reaction can be determined from Equation 5.2. For a negative ∆G◦ value, energy is evolved and the reaction proceeds spontaneously, however when ∆G◦ is positive, energy is absorbed in the process and the reaction will not proceed spontaneously. 5.2 Stability of TiO2 at High-Temperatures 127

Table 5.2: Enthalpy ∆H◦, entropy S◦ and calculated standard free energy ◦ of formation ∆f G values for various compounds. TiO2 (a) = anatase 131, 289–292 and TiO2 (r) = rutile.

◦ ◦ ◦ ∆H at 298 K S at 298 K ∆f G at 1223 K Compound (kJ mol−1) (J mol−1K−1) (kJ mol−1) O2 0.0 205.2 0.0 Si 0.0 18.8 0.0 SiO2 -910.7 41.5 -961.5 TiO -519.7 50.0 -580.9 TiO2 (r) -944.7 50.3 -1006.2 TiO2 (a) -938.7 49.9 -999.7 Ti2O3 -1520.9 77.3 -1615.4 Ti3O5 -2459.4 129.3 -2617.5 TiSi2 -133.9 61.1 -208.6

The entropy and enthalpy data of several relevant compounds are given in Table 5.2. Also included in Table 5.2 are the standard free energies of formation at 950◦C (1223 K), calculated using Equation 5.3. The stability of a TiO2 film in an oxygen ambient is demonstrated by chemical reactions in Table 5.3. In this environment, TiO2 will not react with oxygen to form either Ti2O3 or Ti3O5. The only reaction with a negative standard free energy of formation is the reaction of silicon and oxygen to form SiO2, which proceeds spontaneously at 1223 K.

Table 5.3: The stability of TiO2 in an oxygen ambient is demonstrated by the standard free energy change ∆G◦ at 1223 K for each reaction. The rutile phase of TiO2 was used in the reactions. ∆G◦ Possible TiO2/Si Reactions (kJ) −→ 1 2 TiO2 +O2 Ti2O3 + 2 O2 +395.0 −→ 3 3 TiO2 +O2 Ti3O5 + 2 O2 +398.1

TiO2 +O2 +Si−→ TiO2 + SiO2 -961.5

In a nitrogen ambient, the situation is somewhat different. Several possible furnace reactions are given in Table 5.4, along with the associated standard free energy change ∆G◦ at 1223 K for each reaction. The reactions presented in Tables 5.3 and 5.4 assume that the TiO2 film is rutile. Although all reduction reactions are able to proceed spontaneously, the formation ◦ of Ti2O3 with a ∆G = −165.9 kJ has the most negative free energy change. The strong affinity of the titanium atom for oxygen is evidenced by the large positive standard free energy change in the case of TiSi2. 128 5. Enhancing the Passivation of TiO2-coated Wafers

Table 5.4: Possible reactions between TiO2 and Si in the absence of O2, and the associated standard free energy change ∆G◦ at 1223 K for each reaction. Nitrogen does not play an active role in the reaction and is therefore not shown. ∆G◦ Possible TiO2/Si Reactions (kJ)

2 TiO2 +Si−→ 2 TiO + SiO2 -110.9 4 TiO2 +Si−→ 2Ti2O3 + SiO2 -165.9 6 TiO2 +Si−→ 2Ti3O5 + SiO2 -156.9 TiO2 +2Si−→ TiSi2 +O2 +798.0

Experiment

The centre column of Figure 5.1 shows the processing steps for this “reduction” experiment. Again, n-type 1000 Ω cm FZ (100)-oriented silicon wafers were used. To observe the reduction ◦ of the spray-deposited TiO2 film the samples were loaded at 800 CinanN2 ambient, and the furnace ramped up to 950◦C. After 2 hr the wafers were unloaded from the furnace. The reduction of TiO2 was subsequently avoided by supplying oxygen (O2) to the furnace for the first 10 min. The effect of adding oxygen to the gas ambient will be described in detail in Section 5.3.1.

Results and Discussion

Figure 5.2 compares SEM images of (a) a typical spray-deposited TiO2 film with the familiar texture arising from columnar grain growth, and (b) the region of the N2-annealed TiO2 film that exhibited a bluish-purple hue. The white features in Figure 5.2(b) are precipitates on the surface, while grain boundaries can be seen in top and bottom left-hand corners. The dramatic increase in grains size and the change in surface morphology from Figure 5.2(a) to Figure 5.2(b) would seem to indicate that a sintering process has occurred, resulting in a smoother film with much larger grain sizes (see Figure 5.2(b)). The J0e of the samples with a bluish-purple film was typically very high, of the order of 2 − 4 × 10−11 A/cm2. This is approximately an order of magnitude higher than that achievable with either a bare silicon −12 2 −12 2 wafer (J0e =1× 10 A/cm )oraTiO2-coated wafer (J0e =4× 10 A/cm ). Therefore, the high J0e observed in this experiment could be due to a metallic compound at the TiOx:Si interface, which would explain the high surface recombination velocity. Sub-oxides of TiO2, 58, 172, 173 including TiO and Ti2O3, are known to exhibit metallic conduction properties. The dependence of the conductivity on stoichiometry was observed to change from less than −10 −1 −1 2 −1 −1 141 10 Ω cm for TiO2.00 to 10 Ω cm for TiO1.75.

FTIR Analysis 5.2 Stability of TiO2 at High-Temperatures 129

Figure 5.2: SEM image of (a) a typical spray-deposited TiO2 film, and (b) the region of the N2 annealed TiO2 film that exhibited a bluish-purple hue. The white features are precipitates on the surface, while grain boundaries can be seen in top and bottom left-hand corners.

In order to determine which sub-oxide has formed and to understand the nature of the chemical bonding, FTIR spectroscopic measurements were performed on samples loaded in

O2 and N2. Figures 5.3(a) and (b) compares the spectra obtained from annealing samples in an O2 and N2 ambient, and also a bare silicon wafer reference. In Figure 5.3(a), the absorption peaks at about 260 cm−1,360cm−1, and 430 cm−1 can be assigned to the anatase 261 phase of TiO2. However, Figure 5.3(a) shows that the sample annealed in N2 has a distinctive absorption peak at 493 cm−1, which cannot be assigned to either the anatase or rutile phase of TiO2. The literature discusses two possible origins of an absorption peak for TiO2 samples in this wavelength range. Erkov et al. noted that, firstly, an absorption −1 peak at 470 − 480 cm can be attributed to samples with a thin SiO2 layer at the TiO2:Si interface.117 Secondly, an absorption peak at 480 cm−1 has also been attributed to the 117, 293 presence of Ti2O3.

Figure 5.3(a) shows that the sample loaded in O2 does not exhibit a strong, defined absorp- tion peak at 470 − 480 cm. Additionally, the TiO2 samples placed in the furnace in an O2 ambient are known to possess a thin interfacial SiO2 layer (see Section 5.3.1). Therefore, there is a strong indication that the anomalous absorption peak of the N2 annealed sample is due to the formation of Ti2O3 during high-temperature processing in an N2 ambient. Figure 5.3(b) provides further confirmation of this result, as the weak absorption peaks at 3750 cm−1 −1 294 and 3840 cm can also be attributed to the presence of Ti2O3 on TiO2 substructure.

In summary, spray-deposited TiO2 thin films have been demonstrated to exhibit the following properties when subjected to high-temperature processing. 130 5. Enhancing the Passivation of TiO2-coated Wafers

Figure 5.3: (a) and (b): FTIR spectra of spray-deposited TiO2 films after high-temperature processing in an O2 and N2 ambient.

i) TiO2 films do not result in contamination of the silicon wafer or quartz tube furnace. This was determined by monitoring the bulk minority carrier lifetime of the wafer

before and after TiO2 film deposition and high-temperature processing. Significantly,

a standard commercial 97% pure TiO2 precursor was used for all experiments.

ii) TiO2 films placed in the furnace in an oxygen ambient are stable. After 10 min of O2

the gas flow can be changed to N2 with no adverse effects.

iii) TiO2 films loaded into the furnace in an nitrogen ambient are unstable and the TiO2

is reduced to sub-oxide, most likely titanium sesquioxide Ti2O3.Ti2O3 exhibits very

different optical and electrical properties than TiO2 and is not a suitable AR coating for a solar cell (see Chapter 2).

5.3 Methods of Achieving Surface Passivation with

TiO2 Thin Films

Surface passivation is an extremely important design consideration for high-efficiency solar cells, especially at the front surface where the majority of the light is absorbed. In the previous section, it was demonstrated that spray-deposited TiO2 films are stable under high-temperature processing, as long as oxygen is present at the start of processing. It is widely recognized that a stoichiometric TiO2 thin film, such as those deposited using spray deposition and chemical vapour deposition, on bare silicon affords very little surface 5.3 Methods of Achieving Surface Passivation with TiO2 Thin Films 131 passivation.20, 239, 295, 296 Therefore, this section will investigate the options for capitalising on the required oxidation step in order to enhance the level of surface passivation possessed by TiO2-coated silicon wafers.

A literature review will be presented, and the methods for achieving good surface passivation with TiO2 films will be discussed. The most common method is to grow a thin thermal SiO2 passivation layer on the silicon wafer and subsequently deposit the TiO2 film, however, more recently, limited success has been achieved using non-stoichiometric TiOx films without an SiO2 layer. Following this, experimental results that show that the novel method of growing

SiO2 layers at the TiO2:Si interface after TiO2 film deposition and method of depositing TiO2 on pre-existing phosphorosilicate glass (PSG) layers provides excellent surface passivation.

Deposition of TiO2 on SiO2

The most common method of achieving good surface passivation in conjunction with TiO2

AR coatings has been to initially grow a thin (5−30 nm) thermal SiO2 passivation layer and 16, 20, 22, 283, 295, 297–301 subsequently deposit a TiO2 layer. Thin SiO2 passivation layers have also 150 been used in conjunction with magnesium fluoride (MgF2)/TiO2 double-layer AR coatings.

In that work, no reduction in optical performance was observed for SiO2 thicknesses up to

10 nm. Zhao et al. determined the minimum SiO2 thicknesses for high-efficiency passivated 283 emitter and rear cells (PERC). This was performed by monitoring the Voc and Isc of a textured and planar solar cell as the oxide layer was etched back. The results showed that for a textured surface the Voc started decreasing for a SiO2 thickness less than 25 nm, and at

10 nm for a planar surface. For both textured and planar surfaces the Isc started decreasing for SiO2 thicknesses less than 10 nm.

The performance of BC solar cells fabricated on FZ silicon is limited by the relatively high rear surface recombination velocity exhibited by the alloyed aluminium high-low p-type junction.23 For BC solar cells fabricated on Czochralski (CZ) grade c-Si or multicrystalline silicon (mc-Si) wafers the lower bulk lifetime, typically 20 µs, places additional performance limitations on the device.23 The result is that a poorer front surface recombination velocity caused by using 5 nm-thick SiO2 passivation layers, for example, may not result in a no- ticeable decrease in performance. Additionally, these thinner passivation layers increase the optical transmittance of the TiO2/SiO2 stack and improve device performance. Thin SiO2 layers (5 − 6 nm) have also been demonstrated to provide good surface passivation in some high-efficiency solar cell designs.8 The sensitivity of the surfaces will be somewhat dependent on the doping profile of the emitter. Honsberg et al. published J0e results of TiO2-coated 20 −14 2 emitters with a thin SiO2 layer. An excellent result of 4 × 10 A/cm was achieved. The

TiO2/SiO2 stack has been demonstrated to be compatible with high-efficiency solar cells, 298 with Voc’s as high as 679 mV being achieved in practice. 132 5. Enhancing the Passivation of TiO2-coated Wafers

Growth of SiO2 at the TiO2:Si Interface

Researchers have previously observed that oxygen is capable of diffusing through TiO2 thin 31, 75, 89, 91, 114, 167, 179, 195, 207, 224, 226, 230, 302, 303 films to form a SiO2 layer at the TiO2:Si interface, however none have fully explored the potential until now.35, 50 Previously, the formation of SiO2 has been observed during the fabrication of super-thin capacitors and MOSFETs, 31, 75, 89, 167, 224, 226, 302 which take advantage of the high dielectric constant of TiO2. Brief oxi- dations are performed on such samples in order to reduce the leakage current. However, if the oxidation temperature is too high, the thin SiO2 layer that forms at the interface can be detrimental to the device performance, drastically reducing the effective dielectric con- 31, 197, 230 stant. In the field of PV, three research groups have noted that SiO2 can form at 114, 217 217 the TiO2:Si interface upon heat treatments. In the work of Wong and Waugh, the

TiO2 films were being applied to SP solar cells, which are not able to benefit from improved surface passivation due to the phosphorus “dead-layer” at the top surface. However, it was noted that the films were porous and oxygen could diffuse through the TiO2 film. Szlufcik et al. noted only that the reflectance of the TiO2 AR coating increased after heat treatments at ◦ 900 C, which was attributed to the formation of SiO2 at the interface. Murozono et al. ob- served SiO2 growth underneath TiO2 AR coatings (using Auger electron spectroscopy), when the samples were fired at temperatures up to 1000◦C, but the concept of surface passivation was not investigated.89

The composition of such thin interfacial oxides has also been described in the literature. In the majority of publications related to MOS devices, the interfacial layer is simply assumed 91, 179, 195, 230 to be SiO2 because lower than expected dielectric constants were observed. Auger electron spectroscopy (AES) has been used successfully to identify the interfacial layer as 31, 89 being SiO2. Using RBS, Gartner et al. found that a 9 nm thick SiO2 had formed at the interface after a 1 -hr oxidation at only 300◦C.207 The research performed by Campbell et al. has shown that the dielectric constant of the interfacial oxide is not as low as for pure 75, 167, 224, 226, 302, 303 SiO2. Therefore, it was postulated that the observed 2 − 3 nm thick layer is an amorphous mixed Ti-O-Si oxide. This appears to be likely when considering results that indicate that the very first stages of TiO2 film deposition on a silicon wafer result in 90, 96 an amorphous TiO2 film. Thus, it can be imagined that the initial SiO2 growth at the TiO2:Si interface may mix with the amorphous TiO2 region. Lee observed that non- stoichiometric TiOx oxides at the interface lead to a large leakage current in DRAM cells, and found that it was necessary to include a thin (1 nm) thick SiO2 layer to improve device performance.113

Non-stoichiometric TiOx

Non-stoichiometric TiOx films are typically deposited using physical vapour deposition meth- ods, such as evaporation or sputtering. In this scenario, film growth does not occur due to 5.3 Methods of Achieving Surface Passivation with TiO2 Thin Films 133 a chemical reaction within the system and the oxygen concentration in the ambient plays a critical role in determining the stoichiometry of the film. Three works have demonstrated that some degree of surface passivation can be achieved with non-stoichiometric TiOx thin films deposited directly onto a silicon wafer, however the application of non-stoichiometric films is not pursued in this work.

Wohlgemuth et al. postulated that the higher voltages observed for solar cells with a spray + deposited TiOx AR coating was due to reduced surface recombination velocity at the n Si 49 surface. Crotty et al. demonstrated that solar cells with a TiOx AR coating had a similar open-circuit voltage (Voc) and efficiency (646 mV and 17.5%) to cells that possessed an 8−10 297 nm thick SiO2 passivation layer at the TiOx:Si interface. This method was dependent on the samples undergoing a 400◦C anneal for 5 min in a hydrogen ambient, otherwise the open- circuit voltage and efficiency remained low at about 626 mV and 13%, respectively. More recently, Doeswijk et al. deduced an improvement in surface passivation after depositing 239 TiOx films by laser ablation, based on modulated free carrier absorption measurements. The improved passivation achieved with this film was attributed to fixed positive charges at the TiOx:Si interface, resulting from the oxygen deficiency in the films. The fixed charge reduces the recombination by forming an electric field, which bends the energy bands near the surface of the wafer.

5.3.1 Growth of SiO2 at the TiO2:Si Interface

By initially depositing TiO2 onto the silicon wafer and subsequently performing a brief oxidation, a thin SiO2 layer can be formed at the TiO2:Si interface. The possibility of achieving this result was noted from thermochemistry analysis performed earlier, which ◦ showed that the presence of TiO2 and silicon in an oxygen ambient at 950 C strongly favoured the formation of SiO2 (refer to Table 5.3). The formation of a TiO2/SiO2 passivation stack in this manner, offers several potential advantages over the previously described scheme of depositing TiO2 onto a thermally grown SiO2 layer. These include:

i) The ease of chemical processing, due to the excellent chemical resistance of polycrys- 35, 36 talline TiO2.

ii) The TiO2 film may act as a diffusion barrier to other elements during the high tem- perature processing stage, preventing contaminants from reaching the silicon wafer.

iii) The stoichiometry of the film is ensured as the oxygen (O2) diffuses through the TiO2 layer, removing any oxygen deficiencies.34, 99 This is beneficial for reducing the level of

optical absorption observed in non-stoichiometric TiO2 films deposited by evaporation or sputtering techniques.138 134 5. Enhancing the Passivation of TiO2-coated Wafers

iv) Carbon contamination, resulting from the organo-metallic precursor, is reduced after high temperature processing due to the decomposition of carbonate species.64, 99

v) The refractive index of the film can be tuned by adjusting the deposition and annealing temperatures.148

vi) The method of surface passivation presented here is applicable to all TiO2 films that are able to undergo a brief high-temperature oxidation. To the authors knowledge, this may only exclude highly stressed films deposited by techniques such as PECVD.

An additional advantage of the TiO2/SiO2 stack is that its passivation properties do not 300 degrade under concentrated sunlight, unlike an SiO2 film. This is attributed to the ab- sorption of ultraviolet photons by the TiO2 layer.

Experiment

The aim of this “surface passivation” experiment (see Figure 5.1) was to determine whether good quality surface passivation could be achieved by, first, depositing TiO2 onto silicon wafers, and then subsequently subjecting the wafers to a brief, high-temperature oxidation process.

The wafers used for this experiment were n-type, high-resistivity 1000 Ω cm FZ Si(100). All seven wafers were etched in a sodium hydroxide solution, cleaned in RCA1 and RCA2,194 dilute hydrofluoric acid (HF), rinsed in deionised (DI) water, and blown dry with nitrogen ◦ (N2) before receiving a 175 Ω/2 phosphorus diffusion at 800 C. A lightly-doped emitter was 22 used, firstly, as the J0e will be extremely sensitive to any change in surface passivation. Secondly, similar lightly diffused emitters are used in the fabrication of high-efficiency solar cells.23, 283 Then, the phosphorosilicate glass (PSG) was removed and the wafers well rinsed in DI water and blown dry. Samples 1 − 3 received a spray-deposited polycrystalline TiO2 coating (≈ 70 nm thick) on both sides, with the wafers sitting on a heater block maintained at 450◦C. The processing steps for these samples are shown in Figure 5.1.

Subsequently, the wafers were RCA cleaned again and then loaded into a quartz tube furnace ◦ for the post-TiO2 oxidation. The loading temperature was 800 C and the ambient O2:N2 (1:1). After loading, the temperature was ramped to 950◦C at a rate of 10◦C/min, and the

O2 was switched off after 10 min, leaving a N2 ambient to anneal the samples. The wafers were removed from the furnace at 950◦C after a further 80 min.

As a control, samples 4−7 received an initial oxidation (10 min at 800◦C) after removal of the

PSG. These wafers then underwent TiO2 deposition and oxidation, as per samples 1−3. The thickness of the PSG, and initial and post-TiO2 oxide layers were measured using a G¨artner L116A ellipsometer (λ = 633 nm) and were found to be 30 nm, 5 nm and 7.5 nm, respectively. Transient-PCD measurements were performed in high injection (∆p ≈ 4 × 1015 cm−3) after 5.3 Methods of Achieving Surface Passivation with TiO2 Thin Films 135 each process step in order to determine the emitter saturation current density and to monitor the minority carrier bulk lifetime.282

Results and Discussion

The cross-sectional SEM image in Figure 5.4(a) depicts the as-deposited TiO2 on the silicon substrate, while the growth of a new layer at the TiO2:Si interface can be seen in the

SEM image shown in Figure 5.4(b). The TiO2, which has peeled away from the interfacial

SiO2 layer and substrate during cleaving, seems to be a dense and continuous film, and the interfaces appear to be abrupt. From Figure 5.4(b), the thickness of the TiO2 film and the newly formed interfacial layer are about 67 nm and 6 nm, respectively.

◦ Figure 5.4: SEM images (10 tilt) showing (a) the TiO2:Si interface before oxidation, and (b) the 6 nm interfacial layer grown during the post-TiO2 oxidation (sample 2).

Results of XPS analysis performed at UNSW are shown in Figure 5.5. The top surface of the TiO2 film is at t=0 s, where the Ti:O ratio is about 1:2. A broad SiO2 peak can be seen, centred at t=1600 s. Artifacts of the XPS technique, namely preferential sputtering 304 and islanding, make the SiO2 film appear to be much thicker than is indicated in the SEM image. There is a small amount of carbon adsorbed onto the surface of the TiO2 film, due to the samples being stored in air. The use of an organic precursor may have also resulted in some carbon being retained in the film,89 however the high temperature oxidations are known to reduce these levels significantly.99

The J0e of the wafers after phosphorus diffusion, measured with the PSG on both surfaces, −13 2 was 2.6 × 10 A/cm . The results of transient-PCD measurements after TiO2 deposition and the subsequent oxidation are displayed in Table 5.5. The J0e of all samples exhibits a marked increase after TiO2 deposition due to the poor surface passivation properties of this 136 5. Enhancing the Passivation of TiO2-coated Wafers

Figure 5.5: XPS analysis chemically identifying the interfacial layer as SiO2.

oxide. However, the most significant result is that the J0e of all TiO2 coated wafers is shown −14 2 to decrease to 4.5 − 7.7 × 10 A/cm after the post-TiO2 oxidation. In some cases, the increase in surface passivation has decreased the J0e by nearly two orders of magnitude. The −12 2 poorer passivation (J0e =2.4×10 A/cm ) of samples 4−7 before post-TiO2 oxidation can be attributed to a poor-quality, non-annealed SiO2 layer grown during the initial oxidation step. Notably, this did not limit the final J0e values. The J0e values achieved here with 60 nm thick SiO2 layers compare favourably with the results achieved for high-efficiency, buried-contact solar cells (7.4 × 10−14 A/cm2), with 158 Ω/2 emitters and thermal oxide passivation.22

Table 5.5: Emitter saturation current density (J0e) results measured using transient-PCD. Samples 4 − 7 received an initial oxidation prior to TiO2 deposition, while samples 1 − 3 did not.

2 J0e (A/cm ) Sample After TiO2 After Post-TiO2 Number Deposition Oxidation 1 1.8 × 10−12 5.1 × 10−14 2 1.9 × 10−12 4.7 × 10−14 3 1.9 × 10−12 7.7 × 10−14 4 3.9 × 10−12 6.6 × 10−14 5 4.0 × 10−12 4.5 × 10−14 6 4.1 × 10−12 7.6 × 10−14 7 4.0 × 10−12 5.4 × 10−14

The phosphorus pre-deposition step results in a relatively heavily doped and shallow emit- ter with a surface concentration of 2.0 × 1020 atoms/cm3 and a junction depth of 0.01 µm. 5.3 Methods of Achieving Surface Passivation with TiO2 Thin Films 137

After high-temperature processing the phosphorus surface concentration is reduced to 5.6 × 1018 atoms/cm3 and the junction depth increased to of 0.32 µm. Although some re- distribution of the dopant atoms has occurred during oxidation, the shift in profile cannot fully account for the larger reduction in J0e. Additionally, as the TiO2 films deposited in this work are stoichiometric, the improvement in surface passivation cannot be ascribed to

fixed charges within the films. Thus, the large reduction in J0e observed after the high- temperature oxidation is attributed to the improved surface passivation afforded by the thin interfacial SiO2 layer.

5.3.2 TiO2 on PSG

A simple experiment was performed in order to determine the level of passivation achievable by depositing TiO2 films directly onto the thin phosphorosilicate glass (PSG, P2O5:SiO2) layer remaining after the emitter diffusion step. The use of an existing oxide layer for passivation purposes is of interest for commercially produced solar cells. If successful, it would obviate several steps in a fabrication process, including wet-chemical etching of the PSG layer and the high-temperature growth of an additional passivation layer.

Experiment

Four p-type 1000 Ω cm FZ wafers received a phosphorus emitter diffusion (825◦C for 10 min) from a phosphorus oxychloride (POCl3) source. The POCl3 and O2 concentrations were both 1.5%. This pre-deposition step was followed by a drive-in at 950◦C for 90 min in the same furnace tube. This resulted in a relatively deep emitter with a sheet resistance of 100 Ω/2 and a PSG thickness of 5 nm. TiO2 films (≈ 70 nm thick) were subsequently spray-deposited on both surfaces of wafers 1 and 2 only. These wafers were loaded into an oxidation furnace ◦ at 800 C. The ambient gases were a mixture of both 02 and N2 (3.2 slpm each). The furnace ◦ was ramped to 950 C with the O2 being left on for the first 10 min. The wafers remained in a ◦ N2 ambient for another 80 min before being unloaded at 950 C. Samples 3 and 4 received the same emitter diffusion, however the PSG layer was subsequently removed and a thin (50 nm, sample 3) and thick (1000 nm, sample 4) SiO2 layer was thermally grown. The thin SiO2 layer for sample 3 was grown using the same conditions as the post-TiO2 oxidation above, ◦ while the thick SiO2 layer for sample 4 was grown by loading the sample in dry O2 at 800 C and ramping the furnace up to 950◦C for a total of 90 min. Therefore, all samples received exactly the same high-temperature process, with only the gas ambient being altered. 138 5. Enhancing the Passivation of TiO2-coated Wafers

Results

2 Table 5.6 shows the J0e (in A/cm ) measured using transient-PCD analysis. It can be seen that the J0e increases slightly after TiO2 deposition. The sensitivity of the measurement is due to the minimal thickness of the PSG and the relatively lightly-doped emitter. Finally, it should be noted that after a subsequent oxidation step and anneal, the J0e reduces further by a factor of 2 − 3. There was no evidence of any reaction between the TiO2 film and the PSG layer, before or after high-temperature processing. Table 5.6 also shows the measured

J0e of SiO2 passivated wafers with the same emitter. The passivation achieved with a thin, 5 nm-thick SiO2 layer (sample 3) is very similar to that of the same thickness PSG-layer with a TiO2 coating (sample 2). Improved surface passivation is afforded by the 100 nm- −14 2 thick SiO2 layer with a dark saturation current density of 4.1 × 10 A/cm being achieved.

Table 5.6: Emitter saturation current density (J0e) results measured us- ing transient-PCD. The 5 nm-thick PSG that resulted from the 100 Ω/2 emitter diffusion step was left on the wafer. Only Samples 1 and 2 had TiO2 deposited on them. Samples 3 and 4 had the PSG layer removed and a thin (5 nm, sample 3) and thick (100 nm, sample 4) SiO2 layer was thermally grown.

2 J0e (A/cm ) Sample After POCl3 After TiO2 After Post-TiO2 Number Diffusion Deposition Oxidation 1 1.4 × 10−13 1.8 × 10−13 9.7 × 10−14 2 1.5 × 10−13 2.2 × 10−13 7.7 × 10−14 3 1.3 × 10−13 − 7.5 × 10−14 4 1.3 × 10−13 − 4.1 × 10−14

Discussion

−13 From Table 5.6 it can be seen that a J0e of 1.8×10 can be achieved for TiO2 on P2O5:SiO2 layers, with thicknesses of 70 nm and 5 nm, respectively. This J0e result would limit the Voc 2 of a solar cell with such a 100 Ω/2 emitter to 671 mV (for an assumed Jsc of 33 mA/cm ), which is much greater than a commercially produced BC solar cell (e.g., the best BC solar cell 15 fabricated by BP Solar cell exhibited a Voc of 636 mV ). Hence, such a combined emitter and passivation scheme would not be the limiting factor in the solar cell. The J0e values of wafers passivated with SiO2 only indicate, firstly, that very similar passivation can be achieved with either PSG or SiO2 layers. Secondly, although thicker passivation layers result in a lower

J0e, these films would exhibit greater optical (reflectance) losses than their electrical gains. 5.4 Conclusions 139

One disadvantage of such a passivation scheme is the tight control that is required on the PSG thickness. For the method to be commercially feasible the PSG layer thickness would have to be maintained between 5 − 10 nm. As discussed in Section 5.3, films thinner than 5 − 6 nm have not been successfully implemented in a high-efficiency solar cell design, while thicker films will increase optical losses due to reflection. A thickness of 5 − 10 nm of SiO2 is relatively easy to control by varying the oxidation time and temperature, however the growth rate of P2O5:SiO2 is much higher. A second potential disadvantage of having a PSG layer adjacent to a TiO2 layer is that a reaction between the two films may occur slowly over time. The mechanism for such a reaction is discussed in Chapter 6. The reaction results in a new compound being formed which has both poor optical and passivation properties. No visual signs of a reaction could be observed in these films, however at lower P2O5 concentrations the reaction may still occur over a longer period of time.

5.4 Conclusions

The experiments performed in this chapter have demonstrated that TiO2 is compatible with high-temperature processing without contaminating the wafers or furnaces. The bulk minority carrier lifetimes of samples placed in a furnace for 2 hr at 950◦C were maintained at greater than 2 ms. It was found that the spray-deposited TiO2 films are sensitive to the initial gas ambient in the furnace. Samples loaded in oxygen were stable, however TiO2

films loaded directly into a nitrogen ambient were reduced to a sub-oxide, most likely Ti2O3. The formation sub-oxide was predicted using thermochemistry analysis and confirmed using FTIR spectroscopy.

Other researchers have demonstrated that some level of surface passivation can be achieved using non-stoichiometric TiOx thin films. However, many unanswered questions remain regarding the use of TiOx films, including, firstly, the limits of such a passivation scheme. Secondly, greater knowledge regarding the range of stoichiometries (x values) that afford good surface passivation is required. Thirdly, the exact origin of the surface passivation mechanism needs to be examined, and the role of hydrogen and oxygen vacancies studied. Finally, considering the strong affinity that the titanium atom has for oxygen, the long-term stability of the TiOx passivation scheme needs to be carefully examined.

Importantly, the results of experiments performed in this work have indicated the presence of a thin SiO2 layer, formed by oxidizing the wafer after TiO2 film deposition. This was confirmed using SEM images and XPS analysis. The increase in surface passivation afforded by the interfacial SiO2 layer results in a decrease in J0e by nearly two orders of magnitude −12 2 −14 2 from ∼ 2×10 A/cm after TiO2 deposition to 4.7−7.7×10 A/cm . This demonstrates the ability of the TiO2/SiO2 AR coating to provide excellent surface passivation. The low

J0e and high τbulk values demonstrated here are compatible with high-efficiency solar cells 140 5. Enhancing the Passivation of TiO2-coated Wafers with an open circuit voltage of the order of 700 mV. Any slight reduction in AR coating performance due to the inclusion of the low refractive index SiO2 layer is far outweighed by the passivation benefits.22

−13 2 Finally, it is also demonstrated in this work that a J0e of 1.8 × 10 A/cm can be achieved by depositing a TiO2 film directly onto a 5 nm-thick PSG layer that remains after a 100 Ω/2 2 phosphorus emitter diffusion. For an assumed Jsc of 33 mA/cm ,aJ0e of this magnitude would limit the Voc to 671 mV, which is more than sufficient for commercially produced silicon solar cells. The long term stability of this passivation scheme remains unknown, and possibility of reactions between the PSG and TiO2 layers exist. Chapter 6

Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source

The ability of a dielectric film to act either as a phosphorus dopant source or a phosphorus diffusion barrier would enable a solar cell fabrication sequence to be simplified. In the for- mer case, the phosphorus is incorporated into the antireflection coating and fired at a later time, allowing the P-atoms to diffuse into the silicon. The diffusion barrier is particularly relevant for the buried-contact (BC) solar cell fabrication sequence, where the dielectric film is required to prevent phosphorus from reaching the lightly-doped emitter during the heavy groove diffusion step.

The performance of TiO2 thin films in both of these roles were investigated. It was discov- ered that TiO2 and P2O5 are prone to react chemically at high enough temperatures and concentrations, and that the P-dopant atom significantly alters the properties of the TiO2 film. Thermochemistry analysis and Rutherford backscattering spectroscopy are used to de- termine what chemical reactions are occurring, and the likelihood of success of both of these simplification schemes in the buried-contact solar cell process.

6.1 Introduction

This chapter investigates the performance of TiO2 thin films as a barrier to phosphorus dif- fusion as well as a phosphorus dopant source. If TiO2 could successfully act as a phosphorus diffusion barrier it could most likely be used a direct replacement for the thermally grown silicon dioxide (SiO2) layer implemented in the original buried-contact (BC) solar cell. As well as reducing the number of high-temperature processing steps, this would significant reduce the processing time and cost.

141 1426. Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source

The first two steps in a typical commercial silicon solar cell fabrication sequence are, firstly, emitter formation and, secondly, AR coating deposition. The possibility of combining these two deposition steps by using a phosphorus-doped titanium dioxide (P:TiO2) thin film is also investigated in this chapter. If successful, this would enable the AR coating and emitter dopant source to be deposited in a single low-temperature step, the phosphorus dopants diffusing out from the film and forming an emitter during a subsequent high-temperature process. This could greatly simplify either a solar cell fabrication sequence.

6.2 TiO2 as a Phosphorus Diffusion Barrier

There are several requirements for a stable thin-film diffusion barrier. Firstly, to minimise chemical interdiffusion, the barrier should ideally be a large-grain polycrystalline thin film to reduce grain boundary diffusion.305 Secondly, to prevent chemical interaction the chosen ma- G◦ 305, 306 terial must have a large standard free energy of formation ∆ f . This thermodynamic criterion minimises the chance of a reaction proceeding as described in Equations 5.2 and 5.3. Thus, it is recommended to choose a material such as oxides, nitrides, and transition metal nitrides, borides and silicides.306 Also, the presence of excess nitrogen, carbon, boron, or oxygen in the film can increase its effectiveness as a diffusion barrier.307 It should be noted that diffusion barriers do not eliminate the driving force for diffusion caused by the concentration gradient, but merely reduce rate of interdiffusion.

The majority of the work performed on diffusion barriers has focussed on preventing in- teractions between aluminium and silicon, thus allowing the microelectronics industry to continue to use aluminium contacts under increasingly demanding processing conditions.305

There has been very little work in the literature reporting the use of TiO2 as a barrier to phosphorus diffusion. The only reference to TiO2 acting as a phosphorus diffusion barrier wasmadeinpassingbyDrynanet al.308 In that work, it was observed that when titanium metal films were deposited onto silicon, the titanium reacted with the silicon native oxide to form non-stoichiometric TiOx , and this TiOx layer appeared to limit the diffusion of phosphorus. TiO2 thin films have been demonstrated to act as a diffusion barrier to var- ious other compounds, such as lead titanate,309 and also between copper and ,310 and between platinum-based metal contacts and silicon.311–313 Differing experimental results have been achieved with TiO2 acting as a diffusion barrier to hydrogen (H). Spiegel et al. 314 reported that APCVD-deposited TiO2 films acted as a out-diffusion barrier to hydrogen.

In contrast, Liang et al. found that the sheet resistance of H-doped TiO2 thin films increased from 30 Ω/2 to 500 Ω/2 after two days.58 The increased sheet resistance was attributed to the out-diffusion of hydrogen from the film. Tang et al. postulated that since anatase has a less dense and more defected structure than rutile, this could favour impurity diffusion into the material.120 6.2 TiO2 as a Phosphorus Diffusion Barrier 143

Thermochemistry Analysis

Thermochemistry analysis is used to predict any reactions that may take place between the

TiO2 film and the phosphorus containing furnace ambient. The analysis performed here assumes that the POCl3 and O2 have fully reacted upon entering the neck of the furnace together, forming (P2O5) vapour, which is then deposited as a solid onto the surfaces of the wafer.315 Therefore, the chemical reactions postulated here are between the compounds P2O5 and TiO2. The entropy and enthalpy data of the compounds discussed in this chapter are given in Table 6.1, along with the standard free energies of formation at 950◦C (1223 K), calculated using Equation 5.3. The thermodynamic data is taken from the literature.131, 289–291

Table 6.1: Enthalpy ∆H◦, entropy S◦ and calculated standard free energy ◦ of formation ∆f G values for various compounds. TiO2 (a) = anatase, 131, 289–291 TiO2 (r) = rutile, (g) = gas, (s) = solid. ◦ ◦ ◦ ∆H at 298 K S at 298 K ∆f G at 1223 K Compound (kJ mol−1) (J mol−1K−1) (kJ mol−1)

O2 0.0 205.2 0.0 Si 0.0 18.8 0.0 SiO2 -910.7 41.5 -961.5 TiO2 (r) -944.7 50.3 -1006.2 TiO2 (a) -938.7 49.9 -999.7 (P2O5)2 (g) -3904.1 404.0 -4398.2 (P2O5)2 (s) -3009.9 228.8 -3289.7 TiPO4 -1671.5 96.7 -1789.8 TiP2O7 -2539.7 164.8 -2741.3 TiP3O9 -3255.2 210.9 -3513.1

The six most likely reactions to occur at 950◦C and the standard free energy change ∆G◦ of each reaction, are shown in Table 6.2. Care must be taken when calculating the heats of formation in thin films as some compound phases present in the bulk phase are not observed with thin films.316 Additionally, there is also the consideration of the flux of elements to the reacting surface, as not all elements in the system will be available in the same concentration at the same time to the reacting surface. From Table 6.2 it is noted that TiO2 will react spontaneously with P2O5 to form titanium diphosphate (TiP2O7)or titanium triphosphate (TiP3O9), but the reaction to form titanium phosphate (TiPO4)does not proceed spontaneously at 950◦C. However, if silicon is available to the reaction, noting that silicon and P2O5 are initially present on opposite sides of the TiO2 films, any of the previously mentioned titanium phosphate compounds (TiPxOy) may be formed. 1446. Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source

Table 6.2: Proposed chemical reactions between TiO2 and P2O5 and their standard free energy change ∆G◦ at 1223 K. Proposed Chemical Reaction ∆G◦ (kJ)

TiO2 +P2O5 −→ TiP2O7 -89.8 −→ 1 2 TiO2 +P2O5 2 TiPO4 + 2 O2 +78.5 −→ 1 2 TiO2 +3P2O5 2 TiP3O9 + 2 O2 -78.5

4 TiO2 +Si+2P2O5 −→ 4 TiPO4 + SiO2 -804.6

TiO2 +Si+P2O5 +O2 −→ TiP2O7 + SiO2 -1051.3 1 −→ 2 TiO2 +Si+3P2O5 + 2 O2 2 TiP3O9 + SiO2 -1040.0

6.2.1 Experiment

Several batches of experiments were performed to confirm the results predicted from ther- mochemistry analysis - that some reaction between TiO2 and P2O5 would occur, resulting in the formation of a TiPxOy compound. The substrates used for both diffusion barrier and dopant source experiments were p-type 1000 Ω cm float zone (FZ) (100)-oriented silicon wafers. The diagram in Figure 6.1 depicts the processing steps foe the “diffusion barrier” as well as the “dopant source” experiments. The latter experiment is described in Section 6.3.

Five batches of experiments with eight wafers each were undertaken to determine the effec- tiveness of spray-deposited TiO2 as a phosphorus diffusion barrier. The wafers were loaded ◦ into a POCl3 furnace at 800 C and oxidised for 10 min, before the temperature was ramped ◦ up to 925 CinN2. The POCl3 and O2 concentrations were 1.5% and these gases were left ◦ on for 10 min. A drive-in was performed at 950 C for 90 min in N2 before unloading the wafers. This recipe is typically used to create deep emitters with a sheet resistance of about

45 − 50 Ω/2. Heavier phosphorus diffusions (5 Ω/2) were performed by leaving the POCl3 ◦ and O2 flows on for 90 min at 950 C.

6.2.2 Results

Upon removing the diffusion barrier samples from the POCl3 furnace, it was noted that there were yellow flecks on the surface of the dark blue TiO2 film surface. The experiments were repeated several times with phosphorus diffusions as heavy as 5 Ω/2. For the heavier diffusions it was evident that a reaction between the TiO2 film and the phosphorus had occurred in the furnace, as there were silvery-white, opaque areas on the surface of the film, a significant fraction of the film had a yellow colour to it, while a only a small area remained dark blue. An image of the latter sample is shown in Figure 6.2. In order to determine the nature of the resulting compound, transient-PCD, microwave-PCD, conductivity, and RBS 6.2 TiO2 as a Phosphorus Diffusion Barrier 145

“Diffusion Barrier” “Dopant Source”

NaOH etch (30%, 85°C, 20 min)

RCA2, RCA1, RCA2, HF dip, DI rinse

TiO2 Deposition P:TiO2 Deposition (both sides, (both sides, ~70nm thick, ~70nm thick, Tdep=450°C) Tdep=450°C)

H2SO4:H2O2:DI (1:1:5), RCA2, DI rinse

Furnace Step: Furnace Step: 5–50 Ω/V oxidation (O2) POCl3 Diffusion and anneal (N2)

Figure 6.1: Diagram showing the processes steps used for TiO2 as a phosphorus diffusion barrier and TiO2 as a phosphorus dopant source experiments.

measurements were performed.

Dark blue Opaque region region

Yellow region

Figure 6.2: Image of a TiO2-coated wafer after a heavy (5Ω/2) phos- phorus diffusion. A small area of dark blue coloured film remains, while the majority of the film has a either a yellow or an opaque, silvery white colour. 1466. Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source

PCD and Conductivity Measurements

Transient-PCD analysis was performed in an attempt to determine the effect of the reaction on the surface quality (J0e) and minority carrier bulk lifetime (τbulk). The J0e had increased considerably to 1 − 5 × 10−12A/cm2, primarily due to the poor surface passivation. Mapping of τeff was performed with a microwave-PCD system, however there did not appear to be a relationship to the value of τeff and the white, yellow, or blue regions of the sample. The conductivity of the films was measured using the four-point probe technique. It was found that conductivity increased from the blue region (no measurable conductivity) to the yellow region (able pass a small current through the film) to the white region (about 700 Ω/2). The samples were placed undiluted 48% HF for 24 hours, which slowly etched away the white regions, however the blue regions remained. The remaining blue TiO2 film was removed instantly by subsequently placing the wafer in undiluted NH4OH. The measured conductivity of the diffused layer on the silicon wafer varied from 6000 Ω/2 underneath the white region to > 200000 Ω/2 under the blue region. A hot probe test indicated that an n-type diffusion was present on the surfaces of the p-type wafer. Therefore, it is assumed that a small amount of phosphorus was able to diffuse through the TiO2 film and into the silicon wafer.

RBS Measurements

Rutherford backscattering spectroscopy (RBS) allows the determination of the masses of the elements within a sample, their depth distribution within a 10 nm resolution, and their crystalline structure. Figure 6.3 compares the different RBS spectra obtained from (a) a yellow region, and (b) a silvery-white region of the sample. Curve fitting using the software package RUMP272 suggests that a significant amount of phosphorus has been incorporated into the film. There are two notable differences between the spectra. Firstly, a large amount of silicon has been incorporated into the film in the white region, resulting in the formation of a TiPxSiyOz (see Figure 6.3(b)) compound instead of the TiPxOy compound observed in the yellow region. Secondly, the thickness has dramatically increased from about 70 nm before being placed in the furnace, to 100 nm in yellow regions and 300 nm in the opaque, white regions. A similar phosphorus concentration is observed in both spectra, and the stoichiometry at the surface is similar in both regions.273

6.2.3 Discussion

There is no exact correlation between a likely product from the thermochemistry analysis and the compounds detected using RBS analysis. Whether SiO2 is formed as a reaction by-product or not remains unknown, however it is suspected that the thin SiO2 passivation layer, grown before the POCl3 flow was turned on, has been consumed in the reaction. This, 6.2 TiO2 as a Phosphorus Diffusion Barrier 147

Figure 6.3: Comparison of RBS analysis performed on (a) a yellow region and (b) a silvery-white region of a phosphorus diffusion barrier sample. The solid lines represent experimental data, while the dotted lines are a fit to the experimental data using the simulation package RUMP.272

along with the high J0e values, would seem to indicate that it was not merely a surface reaction that occurred in the furnace and that the entire thickness of the TiO2 film has undergone a reaction, altering the electrical and optical properties of the film.

A search of the relevant literature yields several pieces of information about TiPxOy com- pounds. Ekambaram and Sevov noted that amorphous titanium diphosphate is a white ◦ 317 colour, and that it crystallises in the TiP2O7 phase above 600 C. Glaum and Gruehn observed that it was possible for titanium phosphate compounds to react with silica (SiO2) 290 to form a new titanium silico-phosphate (Ti4P6Si2O25) compound. The ratio between the elements in this new compound are 1 : 1.25 : 0.5 : 6.25, which is remarkably similar to that determined from the RBS data in Figure 6.3(b) for the silvery-white compound (c.f. 1 : 1.25 : 1.05 : 6). The differing fraction of silicon in the compounds formed here could be due to the non-equal availability of reactants to the system. For example, significantly more silicon is available to the reaction from the wafer than the TiO2 or P2O5 constituents. Ad- ditionally, the Si:Ti ratio is too high (1 : 1.05) for the source of the silicon to be solely the 1486. Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source

SiO2 layer, as the ratio of film thicknesses gives about dTiO2 /dSiO2 = 10 : 1. One possible explanation is that the reaction has proceeded throughout the whole film thicknesses and consumed part of the silicon wafer.

The ability of TiO2 to act as a phosphorus diffusion barrier may be better understood by comparing it to SiO2. In comparison to SiO2, the archetypical glass-former, it was noted that the differences between the two compounds included covalent versus ionic bonding, tetrahedral versus octahedral structural units, and that TiO2 is among the worst glass for- 262 mers. Additionally, Zachariasen classified TiO2 as an intermediate or network-modifying oxide, indicating that it does not appear to form a glass by itself.318 The most common network formers in SiO2 are boron and phosphorus, as they are capable of forming glasses 17 by themselves. This allows a P2O5 layer to form a mixed glass when it comes into contact with a SiO2 layer, as occurs in the standard buried-contact process. SiO2 is a more effective diffusion barrier for boron than phosphorous because of the lower temperatures associated 319 with the B2O3:SiO2 glass composition. Examination of the TiO2:P2O5 phase diagram, shown in Figure 6.4, also reveals low temperature glass compositions. However, closer in- spection reveals the formation of the compound TiP2O7 at a wide range of compositions and temperatures. Thus, TiO2 and P2O5 are not miscible, and the two oxides readily react upon contact.

Figure 6.4: TiO2:P2O5 phase diagram, indicating the prevalence for the compound TiP2O7 to form at both a wide range of compositions and at low temperatures.319 6.3 TiO2 as a Phosphorus Dopant Source 149

The prerequisites for the diffusion barrier layer in the BC solar cell fabrication sequence are demanding, requiring that the film remain optically transparent to visible light and maintain excellent surface passivation, in addition to preventing phosphorus from diffusing through to the lightly doped emitter. Although TiO2 prevents all but a small fraction of phosphorus from diffusing into the underlying silicon, the optical, insulating, and passivating properties of the TiO2:SiO2 stack are irreparably damaged in the process, and TiO2 must be regarded as a sacrificial phosphorus diffusion barrier. The author performed one experiment demonstrating that a 200 nm-thick spin-on SiO2 film is sufficient to protect the TiO2 during a heavy groove diffusion, however it was not within the scope of this work to investigate the full potential of suck “workarounds”.

6.3 TiO2 as a Phosphorus Dopant Source

There are several motivations for introducing dopant atoms into thin films. For TiO2,com- mon reasons include altering the electrical conductivity, photoelectric response, chemical re- sistance, optical properties, melting temperature, and the resulting TiO2 crystalline phase, as well as using the film as a dopant source. There are quite a number of references to

P:TiO2 thin films in the literature. In some , the phosphorus was introduced to in- 61, 169, 188 crease the conductivity of a visibly transparent film. In one instance, doping a TiO2 188 film with 1 mol. % of P2O5 increased conductivity of the film by 1000 times. The use of

P:TiO2 in a heterojunction solar cell as an intermediary layer between tin oxide and silicon was proposed, due to its increased electrical conductivity and optimal refractive index.61 181 Moriyama described several methods for depositing P-doped TiOx films. These included e-beam evaporation and sputtering of a phosphorus and titanium source followed by oxida- tion; deposition of a titanium film followed by oxidation of the film in POCl3 atmosphere; deposition of a P-based film onto a TiOx film followed by a high temperature diffusion step; and finally, ion-implantation of phosphorus into a TiOx film. The aim of that work was to suppress the leakage current of a TiOx capacitor.

Another frequent observation is that P:TiO2 enhances the formation of the TiO2 anatase phase, and inhibits the formation of the rutile phase.80, 81, 86 Akhtar et al. deposited doped

TiO2 films formed with TiCl4 and 15% POCl3, and observed an increase in the lattice constant. This indicates that the phosphorus is incorporated into the lattice. Additionally, the smaller phosphorus atoms were observed to diffuse interstitially, which had the effect of inhibiting the phase change to rutile. Rao et al. observed that after 3 hr annealing at 870◦Ca 80 TiO2 film containing 5 at. % phosphorus remained 100% anatase. The enhancement of the anatase phase was also observed when analysing P-doped glass using Raman spectroscopy.86

In the remainder of the works examined here, the phosphorus was intended to be used as a dopant source. Safir initially described use of doped SiO2 as a diffusion source, however 1506. Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source the potential of including dopant atoms, such as phosphorus, into titanium monoxide (TiO),

TiO2, and other thin films was noted. Toyokura and Taguchi demonstrated that after anneal- ing a phosphorus-doped (P:Ta2O5) film on a p-type wafer for 30 min at ◦ 19 −3 320 1000 Ca0.3 µm deep junction is formed with a surface concentration of ND =10 cm . It was mentioned that similar results could be achieved with P:TiO2 however no experimental evidence was given in the paper.

Sharp Corporation (Japan) hold several patents regarding the diffusion of phosphorus from 225, 321, 322 a TiOx antireflection coating into p-type silicon to form a p-n junction. In the work by Yokozawa et al., the titanium and phosphorus precursors, TPT and triethoxy phosphorus, respectively, are introduced into a CVD chamber using an inert gas.225 It is claimed that a subsequent thermal process enables the phosphorus to diffuse from the TiOx film. In subse- quent research, the technique was refined so that highly-doped P:TiOx is deposited where the 321 evaporated contacts will be located. Then, a more lightly-doped P:TiOx layer is deposited over the whole front surface of the p-type wafer. The subsequent high-temperature step then results in the formation of a selective emitter and reduction of the contact resistance. The most recent work by Ui et al., details the variation of the refractive index and resulting sheet 322 resistance with phosphorus concentration of CVD deposited P:TiOx films. By increasing the P:Ti atomic ratio from 0.1 to 1.0 the sheet resistance decreased from 100 Ω/2 to about 30 Ω/2. Two researchers have noted that it is much easier to perform p-type diffusions from boron-doped TiO2 films than n-type diffusions from P:TiO2, however no further explana- tions were given.323, 324 Despite this, Yoldas and Yoldas reported that n-type diffusions could be achieved from TiO2 films containing, firstly, 15% P2O5 and, secondly, 10% triethyl phos- phate. The resulting sheet resistance after high-temperature processing at 1000 − 1050◦Cin 323 N2 for one hour was not stated.

Investigations have also as been performed in the use of titanium silicide (TiSi2) as a phos- phorus dopant source. Privitera et al. achieved shallow junctions (< 100 nm deep) with ◦ P:TiSi2 using thermal treatments in the range 950 − 1150 C. However, this resulted in a high density of silicon phosphide (SiP) precipitates at the interface. La Via et al. noted that titanium-phosphorus compounds were formed in the P:TiSi2, however phosphorus diffusion was still observed due to the similar lattice spacings of the two compounds.325 Another work noted that there was negligible diffusion of phosphorus into the silicon after annealing for 30 min at 800◦C.326 Indium tin oxide has also been doped with phosphorus in order to form an emitter in p-type silicon.327 However, no phosphorus diffused into the silicon, and it is postulated that indium phosphide formed and was evaporated during the subsequent 900◦C thermal treatment.

In this section, initial investigations on the use of spray-deposited TiO2 as a phosphorus dopant source for the formation of emitters in silicon solar cells are reported. The aims were to determine, firstly, whether an emitter dopant could be contained within the dielectric coating, and, secondly, if the dopant atoms could diffuse from the dielectric film into the 6.3 TiO2 as a Phosphorus Dopant Source 151 silicon wafer during a subsequent high-temperature processing. If successful, this would eliminate the need for one high-temperature step, namely the emitter diffusion.

6.3.1 Experiment

In order to realize phosphorus-doped TiO2 films, small volumes (1 wt. % and 5 wt. %) of triethyl phosphate (TEPO, C2H5O3PO, Aldrich, 99%) were added to the TPT, and the solution sprayed onto p-typewafersat450◦C. The samples were then cleaned in boiling RCA2 solution for 5 min and rinsed in DI water. This resulted in the complete etching of the 5% P:TiO2 films, however the 1% P:TiO2 films remained intact. After rinsing and drying, the samples were separated into three batches. All batches were loaded at 900◦C in O2,andaN2 ambient replaced the O2 after the first 5, 10, and 20 min for each of the batches, respectively, while ramping up to 950◦C. In all cases the total processing time was kept constant at 90 min.

6.3.2 Results and Discussion

The sheet resistance of TiO2 films doped with 1 wt. % phosphorus were measured using the four-point probe (FPP) technique. Results indicated that the film itself had a sheet resistance of 300 Ω/2. Transient-PCD analysis indicated that the τbulk of the samples were −13 2 high at 1.4 ms, however the J0e was also high at 9 − 10 × 10 A/cm . After etching off the P:TiO2 film in pure (48%) HF, a hot-probe test confirmed that an n-type diffusion was present, but the conductivity was too low to be measured with the FPP technique. One published work has indicated that titanium atoms are able to diffuse out of a reduced 31 TiOx film, resulting in an n-type diffusion in a p-type silicon wafer. The possibility of this occurring in the spray-deposited samples cannot be ruled out, however the TiO2 films employed here are very close to stoichiometric (2.02 ± 0.13 from RBS measurements) and the strong affinity of titanium for oxygen would seem to rule out any diffusion of titanium into the silicon.

Ferdjani et al. noted that phosphorus can be incorporated into an oxide as oxygenated an- 328 ions PxOy(n-) rather that atomic phosphorus. Additionally, within TiO2, the formation of titanium phosphide (TiP) is common as this compound is very stable.328 XPS experi- ments were performed on the spray-deposited, phosphorus-doped TiO2 films at the Toyota

Technical Institute (Nagoya, Japan). Two 1 wt. % P:TiO2 samples had a total phosphorus concentration at the surface of 3.2 at. % and 2.6 at. %, respectively. The phosphorus 2p peak appears at 135.6 eV while the 2s peak is at 192.5 eV. The height of the phosphorus peaks did not vary significantly with the depth of the film, indicating that the phosphorus dopant 87 atoms are quite uniformly incorporated into the TiO2 lattice. 1526. Performance of TiO2 Thin Films as a Phosphorus Diffusion Barrier and Dopant Source

It would seem that for P:TiO2 films to act as an effective dopant source, phosphorus dopant concentrations of about 15% would be required along with higher processing temperatures, 323 as reported by Yoldas. However, if a 1 mol. % phosphorus-doped TiO2 film increased the conductivity by 1000 times,188 a 15% concentration would most likely make the films very conductive.

6.4 Conclusions

The results from this work indicate that a 70 nm thick TiO2 film can function either as a phosphorus diffusion barrier or a phosphorus dopant source. Although the TiO2 diffusion barrier resulted in only very light phosphorus diffusions underneath the film, a reaction with

P2O5 substantially alters the optical and electrical properties and limiting its usefulness.

Therefore, TiO2 must be described as a sacrificial diffusion barrier to phosphorus or, alter- natively, the TiO2 must be protected by another film (such as SiO2) during the diffusion process.

The ability of TiO2 to act as a dopant source for phosphorus atoms is also somewhat limited.

The increased conductivity of the TiO2 film due to the phosphorus incorporation is unlikely to be compatible with the insulating requirements of the BC solar cell electroless metal plating step. Further investigations would need to be performed to ascertain the benefit of

P:TiO2 films in other solar cell structures, such as those using screen-printed and evaporated metal contacts. However, it is anticipated that the increased optical absorption in the films, along with the lengthy diffusion times and temperatures involved (> 1hrat > 1000◦C) will limit the industrial uptake of this technology, even for screen-printed solar cells. Chapter 7

TiO2 Antireflection Coatings

7.1 Introduction

Commercial solar cells fabricated on multicrystalline silicon (mc-Si) wafers exhibit poorer short current densities (Jsc) than their monocrystalline counterparts, primarily due to the inability to adequately texture mc-Si wafers. This is due to the randomly oriented grains in mc-Si material, and after texturing in a basic (NaOH or KOH) solution, only 20 − 30% of the grains result in the familiar {111} pyramid faces. The pyramids enable the majority of light striking the surface to have two chances at being transmitted into the silicon, instead of being reflected. Research has been performed on alternate texturing methods, such as reactive ion etching,329–331 laser- and photolithographically-defined texturing,28, 332 mechan- ical texturing333, 334 and acid-based chemical etching to form either textured335 or porous silicon.336, 337 While some of these methods enable low reflectance to be achieved on mc-Si wafers, none of these techniques would appear to be particularly attractive in a commercial environment, due to either the high costs (photolithography, maintaining a vacuum) or large quantities of dangerous chemicals involved.

An alternate method of reducing the front surface reflection is with antireflection (AR) coatings. This chapter will briefly review the theory of AR coatings, before establishing the performance of existing silicon nitride and TiO2 AR coatings. Subsequently, a novel and simplified approach for the formation of TiO2 single- and multi-layer antireflection AR will be presented and the performance of these coatings established.

Figure 7.1 shows the radiation intensity (φ) of the AM1.5 global spectrum (in mW/cm2)as a function of wavelength. The data is taken from Green and is normalised to 100 mW/cm2.8

Also shown in Figure 7.1 is the available short-circuit current density (Jsc)ateachwave- length. An AR coating should exhibit a minimum reflectance at about 600 nm in order to take full advantage of the peak in the spectrum at 550 − 750 nm. A maximum Jsc of 45.97 mA/cm2 can be obtained in the 300 − 1200 nm wavelength region.8

153 154 7. TiO2 Antireflection Coatings

Figure 7.1: Global AM1.5 spectrum and short-circuit current density for wavelengths 300 − 1200 nm. The data was taken from Green.8

There are two common standards for measuring the performance of a texturing scheme or

AR coating. The first is the weighted average reflection Rw,definedas: λmax R(λ)N (λ)d(λ) λmin ph Rw = .(7.1) λmax N (λ)d(λ) λmin ph

In Equation 7.1, the values used for λmin and λmax are typically in the range 300 − 400 nm and 1000 − 1200 nm, respectively. The amount of solar flux under the AM 1.5 spectrum is represented by Nph(λ), while Rw is the reflectance of the solar cell. A Rw that is evaluated over a narrower wavelength range is likely to be lower than that evaluated over a wider wavelength range. This is especially true if the short wavelength regions (< 400 nm) are excluded, where the reflectance of silicon is very high, or if the long wavelengths (> 1000 nm) are excluded, where reflectance from the back surface of the silicon wafer is observed. The advantage of using Rw as a figure-of-merit is that it is purely optical and does not depend on the electrical performance of the solar cell.

The second measure of performance is the short-circuit current density Jsc, as shown in Equation 7.2, λmax Jsc = q (1 − R(λ)) Nph(λ)IQE(λ)d(λ), (7.2) λmin where the electrical response of the solar cell as a function of wavelength is contained in the internal quantum efficiency (IQE)term,Nph is the photon flux of the solar spectrum, and q is the electronic charge with the value 1.602×10−19 C. For comparing texturing methods and 7.2 Previous Developments in AR Coatings 155

the majority of AR coatings, the quantity Rw will be used as many of these methods could be applied to either screen-printed or high-efficiency solar cell designs. In later sections, the

Jsc will be used to determine the performance of single- and multi-layer TiO2 AR coatings for buried-contact (BC) solar cells. The advantage of using Jsc is that this is one of the important final output parameters of a solar cell.

7.2 Previous Developments in AR Coatings

This section will briefly introduce the theory behind single- and double-layer AR coatings, before reviewing the two most common AR coatings in the PV industry, TiO2 and silicon nitride (SiNx). For really high-efficiency, laboratory-scale solar cells, the evaporation of materials such as zinc sulphide (ZnS) and magnesium fluoride (MgF2) is a standard process, however this technology has no relevance in a production setting due to the high-cost and low-throughput nature of the evaporation method.

7.2.1 Theory and Design of AR Coatings

Single Layer AR Coatings

A single layer antireflection (SLAR) coating is the minimum requirement for any silicon solar cell produced today. Light is well absorbed by semiconducting materials such as silicon, however these materials exhibit high refractive indices. For example, silicon has a refractive 8 index of nsi =3.939 at 600 nm. This refractive index is much greater than air, which has a constant refractive index of n0 =1.0, and glass (n0 =1.52 at 600 nm). The reflectance of normally incident light at such an interface is given by 2 n − n0 R = Si ,(7.3) nSi + n0 which means that 35.4% or 19.6% of the light is reflected off an air:silicon or glass:silicon interface in the first bounce, respectively. If an optimum-thickness AR coating is inserted between the silicon and ambient medium, the minimum reflectance is given by 2 2 n − n0 n R = AR Si , (7.4) n2 n n AR + 0 Si where nAR is the refractive index of the coating. To achieve zero reflectance at one wave- length, the value of nAR should be √ nAR = n0 nSi (7.5) and the film thickness (dAR) must meet the quarterwave optical thickness requirement

λ0 dAR = , (7.6) 4 nAR 156 7. TiO2 Antireflection Coatings

where λ0 is the wavelength of zero or minimum reflectivity.

This indicates that an AR coating for a silicon solar cell in air should have a refractive index of 1.985 and a thickness of 75.6 nm, while a glass encapsulated cell requires an AR coating with nAR =2.447 and dAR =61.3 nm. Figure 7.2 shows the calculated reflectance for four different SLAR coatings. The lines in Figure 7.2 represent theoretical AR coatings with fixed (non-dispersive) refractive indices of 1.985 and 2.447 (as determined above), while the data points are two TiO2 films deposited by the author, indicating that an excellent match between the theoretical and practical coatings can be achieved. The reflectance minimum of the encapsulated silicon is broader due to the higher refractive index of the glass, however the first reflectance bounce off the air:glass interface means that only a minimum of 4.3% can be achieved at the design wavelength of λ0 = 600 nm. The weighted average reflectances

(350 − 1150 nm) of the fixed AR coatings are Rw =8.75% and Rw =8.41% for air and glass encapsulation, respectively. In comparison, the minimum weighted average reflectances that can be achieved using experimental TiO2 coatings are Rw =9.45% and Rw =8.98% for air and glass encapsulation, respectively. The TiO2 AR coatings deposited in this work will be discussed in detail in Section 7.4.

Figure 7.2: Modelled reflectance of the fixed refractive index coatings, and two TiO2 AR coatings deposited in this work. The performance of the TiO2 coatings is extremely close to that of the theoretical coatings.

Double Layer AR Coatings

Whereas a SLAR coating can be designed to achieve zero reflectance at one wavelength, a double layer antireflection (DLAR) coating offers a further reduction in reflectance. In this 7.2 Previous Developments in AR Coatings 157

scenario, the refractive indices are stacked as follows nSi >nAR2 >nAR1 >n0, where nAR2 and nAR1 represent AR coatings with high and low refractive indices, respectively. For layers with equal optical thickness, such that nAR2 dAR2 = nAR1 dAR1 = λ0/4, the reflectance at λ0 becomes 2 2 2 n n − n n0 R = AR1 Si AR2 .(7.7) n2 n n2 n AR1 Si + AR2 0

The reflectance R will be either a maximum or minimum at λ0 depending on the relationship n2 n n2 n between the refractive indices. From Equation 7.7, it can be seen that if AR1 Si = AR2 0, zero reflectance at λ0 will be achieved. This is a much broader minimum than is achievable with a SLAR coating. If nAR1 nAR2 = n0 nSi a maximum of R will occur at λ0 and double zero reflectance will be achieved at two wavelengths either side of λ0. The latter configuration is commonly used to achieve minimum reflection across the whole solar spectrum (300 − 1200 nm).

For two quarter-wavelength coatings, the optimal refractive indices of each layer in a DLAR stack can be determined by279

n3 n2 n n3 n n2 . AR1 = 0 Si and AR2 = 0 Si (7.8)

The mathematics for determining the reflectance off double- or multi-layer AR coatings is not difficult, however it is much more laborious than the single coating case above, especially for films that exhibit absorption. The interested reader is referred to texts that discuss this optical theory such as Heavens38 and Macleod.279

Absorbing AR Coatings

In many instances, the amount of absorption in the AR coating is negligible338 and Equation

7.1 (see Section 7.1) can be safely applied in order to determine Rw. However, for AR coatings possessing high refractive indices, absorption is commonly observed in the short wavelength spectrum35, 339, 340 and the use of Equation 7.1 will lead to an erroneous answer.

This problem can be solved by considering the weighted average transmission Tw through the AR coating, defined as λmax T (λ)N (λ)d(λ) λmin ph Tw = .(7.9) λmax N (λ)d(λ) λmin ph

This performance parameter is somewhat more complicated to determine experimentally, compared to Rw. In order to determine Rw, the reflectance of the sample is simply mea- sured using a spectrophotometer with an integrating sphere, and a numerical integration is performed using a mathematical software package, such as Mathematica, giving a value for

Rw. In order to determine Tw, there are several options. Firstly, careful optical modelling 158 7. TiO2 Antireflection Coatings of the reflectance data can yield information regarding the extinction coefficient of the film. A certain amount of prior knowledge is required to ensure that a “physically real” answer is reached. Secondly, this analysis can be greatly assisted by depositing the AR coating in question onto a transparent substrate and measuring the transmittance. This provides valuable data to assist in the determination of n and k. The only drawback to this approach is that films such as TiO2 may have a different preferential phase when deposited onto glass or quartz compared to silicon (see Section 2.2.3), and silicon nitride films deposited using PECVD may possess different optical properties due to the insulating nature of the glass substrate. Thirdly, an alternate method for obtaining the optical constants can be used, such as ellipsometry. Once n and k are known the transmittance or reflectance curve can be generated using software and Tw or Rw calculated.

The remaining quantity, average weighted absorptance Aw, then becomes another useful

figure of merit for solar cell design. The Aw is defined as λmax A(λ)N (λ)d(λ) λmin ph Aw = .(7.10) λmax N (λ)d(λ) λmin ph The inclusion of absorption into the optical model also means that the equation for deter- mining the Jsc of a solar cell, previously defined in Equation 7.2, must be modified and now appears as λmax Jsc = q (1 − R(λ) − A(λ)) Nph(λ)IQE(λ)d(λ) λmin λmax = q T (λ)Nph(λ)IQE(λ)d(λ). (7.11) λmin

The wavelength limits used in this work are λmin = 350 nm and λmax = 1150 nm. These limits were imposed by the spectroscopic ellipsometer, for which the data outside the 350 − 1150 nm range is typically quite noisy. The solar spectrum for silicon solar cells is usually 300 − 1200 nm, however only a small fraction of light is emitted from the sun at λ<350 nm (≈ 0.35 mA/cm2) and is absorbed by silicon at λ>1150 nm.

The electrical response of a silicon solar cell varies with wavelength due to the varying absorption coefficient of silicon, as well as the level of front and rear surface passivation and the quality of the silicon substrate. Silicon wafers with a lower bulk minority carrier lifetime

τbulk typically exhibit a reduced response to long wavelength light due to defects within the crystal structure. This includes c-Si wafers grown with the Czochralski (CZ) method, which have a low τbulk due to a high oxygen and carbon concentration, and mc-Si wafers, where the grain boundaries between the crystallites reduces τbulk.

In Equation 7.11, the IQE(λ) term indicates the electrical response of the solar cell over the solar spectrum, called internal quantum efficiency (IQE). The IQE of the solar cell is strongly influenced by the quality of surface passivation, primarily at the front and to some degree 7.2 Previous Developments in AR Coatings 159

at the rear. The use of thin SiO2 layers at the AR coating:silicon interface was discussed in Section 5.3. The IQE of a simulated BC solar cell, used to predict the performance of solar cells with TiO2 AR coatings, is plotted in Figure 7.26 in Section 7.5.

Materials for AR Coatings

Common materials used for AR coatings are summarised in Table 7.1. The refractive indices may vary somewhat depending upon the deposition method, deposition rate, and any post- deposition annealing steps. Transparent conducting oxides such as indium tin oxide and , commonly used in thin film devices, are not shown as their electrical conductivity severely compromises their transparency.

Table 7.1: Refractive indices of materials commonly used in antireflection coatings (adapted from Green7).

Material n (λ = 600 nm)

Magnesium fluoride MgF2 1.38 Silicon dioxide SiO2 1.46 Aluminium dioxide Al2O3 1.8 Silicon monoxide SiO 1.8 − 1.9 dioxide CeO2 2.2 Tantalum pentoxide Ta2O5 2.1 − 2.3 Silicon nitride SiNx 1.9 − 2.1 Titanium dioxide TiO2 1.9 − 2.4 Zinc sulphide ZnS 2.3 − 2.4

Revesz made several important observations about the desirable properties of an AR coat- ing material.43 Firstly, the use of a non-crystalline material is important in AR coating design because the grain boundaries in polycrystalline film cause scattering and decrease transparency. Non-crystalline materials have also proven to be most suitable for fabricating capacitors. However, it is desirable for the noncrystalline material to exhibit a reasonable degree of short range order in order to avoid increased absorption due to unsaturated bonds. Thus, a vitreous material is deemed to be more suitable than an amorphous material. A good example of a vitreous material is thermally grown SiO2, and this would be expected to exhibit lower absorption than sputtered amorphous SiOx, for example. TiO2 films grown in a chemical process at low temperatures, such as spray-deposition or CVD, could also be expected to exhibit a high degree of short-range order.

Roger and Colardelle noted that as a common rule for pre-determining the optical properties of a thin film: a slow deposition rate leads to low absorption and low refractive index, whereas a fast deposition rate is conducive to a high refractive index, but also high absorption.277 160 7. TiO2 Antireflection Coatings

Thus, when designing multilayer AR coatings with absorbing materials, there may be a design trade-off with perfectly matching the refractive indices or avoiding high absorption.

Optical Modelling Software

Two pieces of optical modelling software were used extensively throughout this thesis. Firstly, WVASE,62 already mentioned in Section 4.6, was used for analysing measured ellip- sometry and spectrophotometry data. This software was chosen for its speed, robustness, and a large number of integrated models. A double Lorentz oscillator model was success- fully used for fitting the Ψ and ∆ data and extracting the dispersive refractive indices n and extinction coefficients k in the range 350 − 1150 nm. The n and k values were exported from WVASE into TFCalc341 for optimisation of the AR coating layer thickness. TFCalc enables the user to set either distinct or a range of wavelengths over which to minimise the reflectance. A simplex downhill algorithm was normally used for optimising the layer thick- nesses, and a passivating oxide layer could also be included at the interface. The reflectance, transmittance, and absorptance data was then exported to a Mathematica program where the average weighted reflectance, transmittance, and absorptance values (see Section 7.2.1) were calculated, along with the maximum short circuit current density, Jsc.

7.2.2 TiO2 AR Coatings

16, 19 The most commonly used material in the PV industry for AR coatings is TiO2. These are SLAR coatings, primarily deposited onto screen-printed solar cells. This is primarily due to the excellent optical properties of TiO2, its ease of deposition and low cost. Double- layer antireflection (DLAR) coatings are used in many high-efficiency solar cells, in order to further reduce front surface reflectance losses. These DLAR coatings are normally utilised in laboratory-scale devices, and the coatings are optimised for measurements in air. However, some and, secondly, under glass encapsulation. Ultimately, however, any successful solar cell design will require encapsulation to protect the device from the environment, and a review of DLAR coatings under glass will be provided.

SLAR Coatings

The review paper by Kern and Tracy provides an excellent review of early AR coating 46 methods and development, focussing particularly on TiO2. In that work, screen-printed solar cells exhibited a 41% increase in efficiency after receiving an optimised TiO2 AR coating. Many other researchers have observed a 42 − 50% increase in efficiency after depositing a 89, 157 TiO2 AR coating, the higher numbers coming from glass-encapsulated coatings with 89 nAR =2.4. 7.2 Previous Developments in AR Coatings 161

Brinker and Harrington optimised the ratio of SiO2 to TiO2 in similar SLAR coatings, and 201 found that a 90% TiO2/10% SiO2 mixture achieved the lowest reflectance. The bandwidth, or range of wavelengths, at which the reflectance was ≤ 5% was also determined for various

AR coatings, with the 90% TiO2/10% SiO2 SLAR achieving 196 nm. This value was 10% higher than the SLAR coating of Yoldas and O’Keeffe above, and it was concluded that this was due to a slight gradient in the refractive index.

DLAR Coatings

Figure 7.3 shows a plot of the maximum Jsc achievable for a non-encapsulated silicon solar 45 with optimised single-, double-, and triple-layer AR coatings. The Jsc increases from 25.7 mA/cm2 with no AR coating, to 36.4 mA/cm2 with a SLAR coating. A DLAR coating 2 improves the Jsc further to 38.7 mA/cm , while a triple layer antireflection (TLAR) coating 2 can theoretically achieve a Jsc of 39.2 mA/cm . Thus, a theoretical increase in current of 6 − 8% provides the motivation for researchers to develop multi-layer antireflection (MLAR) coatings for the photovoltaics industry. In reality, a TLAR coating provides little benefit over a DLAR coating due to the limited range of transparent materials that are available. Currently, DLAR coatings are only used in industry for small-scale production of high- efficiency solar cells.301 This is primarily due to different materials and different deposition systems being required for each of the DLAR coating layers. Zhao and Green noted that encapsulated DLAR coating designs using high refractive indices (nAR2 > 2.5) do not improve much above the SLAR coating due to absorption in the lower coating.338 An additional advantage of using a DLAR coating is the reduced sensitivity to layer thicknesses,342, 343 while a disadvantage is that the performance of a DLAR coating decreases rapidly once a 344 SiO2 passivation layer is included at the silicon interface.

Non-Encapsulated Solar Cells

A number of researchers that high refractive index TiO2 films are suitable for being paired up with low refractive index SiO2 films to obtain excellent performance from non-encapsulated DLAR coatings.44, 97, 150, 152, 342, 344–348 In a study of many possible AR coating materials, Jelli- 344 son and Wood determined that a SiO2/TiO2 DLAR coating had the highest performance.

Martinet et al. also concluded that SiO2/TiO2 DLAR coatings closely approached the ideal, with optimal refractive indices of nAR1 =1.58 and nAR2 =2.50 for a silicon in air system at 97 λ0 = 600 nm. Swartz et al. used SiO2/TiO2 DLAR coatings for silicon concentrator solar cells, operating up to 350 suns.345 Yoldas developed a DLAR coating comprised of a top layer of 10% TiO2/90% SiO2 (nAR1 =1.4)andabottomlayerofTiO2 (nAR2 =2.4), which increased the performance of uncoated silicon solar cells by 49%.346

Bouhafs et al. demonstrated that a Rw of 2.65% (300 − 1100 nm) could be achieved for 347 a SiO2/TiO2 DLAR coating, slightly less than the modelled result of 2.39%. Modelling results from Wang et al. indicated that a SiO2/TiO2 DLAR coating in air should be able 162 7. TiO2 Antireflection Coatings

Figure 7.3: Maximum Jsc achievable for a non-encapsulated silicon solar cell with optimised single-, double-, and triple-layer AR coatings. The films used are hypothetical and absorption is assumed to be zero (adapted from Nagel et al.45).

2 44 to achieve a Jsc of 42.2 mA/cm under AM0 illumination. Jiao and Anderson deposited 152 SiO2/TiO2 DLAR coatings on metal/insulator/n-Si/p-Si (MINP) solar cells. The Jsc and efficiency of the solar cells under AM 0 illumination improved by up to 46% after receiving 2 the DLAR coating, with a maximum Jsc of 44.3 mA/cm being observed. No degradation in Voc or fill factor (FF) was observed after evaporating the SiO2 and TiO2 coatings, as has been observed by other researchers.348

342 Pettit et al. used spin-on SiO2 and TiO2 films as a DLAR coating. Aweightedaverage reflectance, including reflectance from the rear silicon:air interface, of 5% was achieved for experimental films, while the minimum achievable Rw with these films was modelled as being 3.8% over the wavelength range 320 − 1120 nm. Figure 7.4 shows the Rw contours for a wide range of SiO2 and TiO2 layer thicknesses. It can be seen that a wide range of layer thicknesses, ±15% for SiO2 and ±10% for TiO2, can still achieve a Rw of 4.3%, just 0.5% above the minimum. It was also noted that these results were better than the Rw of 7% 349 achieved by Johnson for a SiNx quadruple-layer AR coating (see Section 7.2.3).

Kamataki et al. performed modelling to determine the optimum SiO2 and TiO2 layer thick- 150 nesses in a MgF2/TiO2 DLAR coating with a SiO2 passivation layer. Figure 7.5(a) shows the effect of increasing the SiO2 passivation layer thickness and the optimum TiO2 film 2 thickness. The MgF2 thickness was fixed at 100 nm. A theoretical Jsc of 42.6 mA/cm was achieved on a planar surface for SiO2 and TiO2 thicknesses of 6 nm and 60 nm, respectively. 2 For a V-grooved surface, the potential Jsc increased to 44.1 mA/cm . Modelling also demon- strated that the optical performance of the DLAR coating did not change significantly for

SiO2 thicknesses less than 10 nm, as shown in Figure 7.5(b). 7.2 Previous Developments in AR Coatings 163

Figure 7.4: Optimised layer thicknesses of a SiO2/TiO2 DLAR coating with refractive indices nAR1 =1.414 and nAR2 =2.243 (at λ = 632.8 nm), respectively. The ’X’ marks the experimental DLAR coating with Rw = 5% (including back reflectance), while the minimum Rw for this coating 342, 343 is shown at •, corresponding to 95 nm SiO2 and 62 nm TiO2.

Modelling performed by Nagel et al. indicated that a PECVD-deposited SiNx film could 45 be used a surface passivation layer under a SiO2/TiO2 DLAR coating. The SiNx has a significantly higher refractive index than SiO2 (2.2 vs. 1.46), reducing the reflectance losses 2 due to the inclusion of a passivation layer. A Jsc of 37.7 mA/cm was achieved with this DLAR coating, however the use of several different materials in one DLAR coating is likely to make this coating too costly in a production environment.

Encapsulated Solar Cells

Research performed at Solarex found that hot-sprayed DLAR Al2O3/TiO2 DLAR coatings exhibited a 10% performance advantage over TiO2 SLAR coatings, and further enhancement was achieved once a cover slide was in place.49 The hot-sprayed coatings exhibited high refractive indices, ideal for DLAR coatings under glass. In a subsequent study, it was found that incorporating the Al2O3/TiO2 DLAR coating in a screen-printed solar cell process would result in an 0.4% (absolute) increase in efficiency, reducing the cost of the module by 350 US$0.03/Wp (1990 dollars).

Wang et al. performed experiments with many SLAR and modelled DLAR coatings in air and under fluorinated propylene (FEP) plastic sheets under AM0 illumination.44 2 The best results of Jsc =44.0 mA/cm were achieved using a SiO (63 nm)/TiO2 (45 nm) DLAR coating under the FEP (n0 =1.34) cover.

Herpin equivalent layers can be used to create a layer with a desired refractive index, when a single material with that refractive index is not available. Aiken used Herpin equivalent layers to fabricate a quadruple-layer antireflection (QLAR) coating using TiO2 and Al2O3 stacked in 164 7. TiO2 Antireflection Coatings

Figure 7.5: (a) Variation of calculated Jsc as a function of TiO2 film thickness and SiO2 passivation layer thickness for V-grooved and planar silicon surfaces, and (b) optical dependence of Jsc on SiO2 passivation layer thickness. Note that the change in electrical passivation with SiO2 is not included in the purely optical model (adapted from Kamataki et al.150).

a high-low-high-low (HLHL) pattern.351 Modelled results have shown that a QLAR coating can have a solar weighted reflectance as low as 7.0% (wavelength range not specified), nearly matching the optical performance of a TLAR step-down (i.e., nAR1

Cudzinovic et al. deposited Al2O3/TiO2 and SiO/TiO2 stacks by e-beam and thermal evap- oration, respectively, as DLAR coatings for high-efficiency solar cells.301 Severe degradation of the surface passivation quality was reported for e-beam evaporation, while the damage caused by thermal evaporation was also significant. This damage was most apparent for textured samples, and only a fraction of the surface passivation could be recovered by per- forming a forming gas anneal (FGA). 7.2 Previous Developments in AR Coatings 165

Textured AR Coatings

Gee et al. and Liang developed several textured SLAR coatings for application in mc-Si solar 58, 218, 352 cells. While experiments were performed with APCVD-deposited ZnO and TiO2,the best results were obtained using CVD-deposited diamond films (nAR1 ≈ 2.4). The weighted average reflectance of the encapsulated sample was 7.3% over the wavelength range 300 − 218 1000 nm. The TiO2 films exhibited poor adherence and the research was discontinued.

The best results were achieved using an evaporated TiO2 layer covered with a textured ZnO layer. Even with the lower refractive indices of the ZnO film (nAR1 ≈ 2.0), the excellent 218 texture reduced the Rw to 6% (300 − 1000 nm).

7.2.3 Silicon Nitride AR Coatings

SLAR Coatings

Silicon nitride films deposited by plasma enhanced chemical vapour deposition (PECVD) have recently become popular due to their large hydrogen fraction (about 25%), which can provide some level of surface passivation and assists in increasing the bulk minority carrier lifetime in mc-Si wafers. The disadvantage to these coatings is the high cost of depositions due to maintaining a vacuum.

Nagel et al. compared SiNx and TiO2 SLAR coatings for encapsulated and non-encapsulated silicon solar cells.45 Their results show that while excellent performance can be achieved using

SiNx films in air, the glass-encapsulated results show that the performance of the SiNx SLAR 2 2 (Jsc =34.1 mA/cm ) and TiOx/20 nm-SiO2 (Jsc =34.0 mA/cm ) coatings are similar once surface passivation is taken into account. As discussed in Section 5.3.1, 10 nm of SiO2 is usually sufficient to act as a good surface passivation layer, and if included in the modelling this would increase the Jsc of the TiO2 coatings slightly, due to decreased reflectance losses incurred by the low refractive index SiO2 layer.

Doshi et al. performed an extensive study into AR coating optimisation using PECVD silicon nitride films.339 Initially, ten films with different refractive indices were deposited, ranging from 2.03 to 2.55 (at λ = 632.8 nm. The photocurrent lost due to light reflected off the front surface or absorbed in the AR coating was defined as λmax Jpcl = q R(λ)Nph(λ)IQE(λ)d(λ)+ λmin λmax q A(λ)Nph(λ)IQE(λ)d(λ)(7.12) λmin = Jpcl(refl) + Jpcl(abs) . (7.13)

Minimizing the Jpcl is identical to maximising the Jsc. A 100% IQE in the range 400−1100 nm 166 7. TiO2 Antireflection Coatings

2 339 was assumed with 0% IQE elsewhere, for a maximum Jsc of 41.5 ,mA/cm . The highest performance SiNx SLAR coating under glass was achieved with nAR1 =2.23 (λ = 632.8 nm), 2 with Jpcl =2.24 mA/cm , Rw =4.37%, Aw =1.03% and Tw =94.60%. This value is significantly lower than the ideal value of nAR1 =2.45 calculated previously (see Section

7.2.1). Doshi et al. determined that while the reflected losses were lower for a SiNx film with nAR1 =2.42, the increased absorption losses (Aw =2.66%) reduced the overall performance of that SLAR coating. Performing a similar optimisation for a SiNx SLAR coating in air showed that a film with the lowest refractive index nAR1 =2.03 resulted in the lowest losses, 2 with Jpcl =3.32 mA/cm , Rw =7.98%, Aw =0.02% and Tw =92.00%.

BP Solar have replaced rudimentary silicon dioxide (SiO2) AR coating, originally imple- mented in the BC solar cell fabrication sequence, with a low-pressure chemical vapour de- 15, 16 position (LPCVD) silicon nitride (SiNx) layer. LPCVD SiNx films typically have a 17 2 refractive index of about 2.3 at 600 nm. These films have achieved a Jsc of 38.3 mA/cm (unconfirmed) on the best BC solar cell fabricated on a textured crystalline silicon wafer.15

DLAR Coatings

Two solutions to minimizing the number of materials in a DLAR coating have been reported, 339, 340 both using SiNx films. Winderbaum et al. used SiNx films with refractive indices of nAR1 =1.95 and nAR2 =2.5 as a DLAR coating for passivated emitter solar cells (PESC) 340 and SP solar cells. Figure 7.6 compares the optical properties of the spray-deposited TiO2 from Solarex and the bottom layer SiNx film from the DLAR coating. The performance of the TiO2 is superior, exhibiting a much lower absorption while maintaining a higher refractive index. When applied to a PESC cell, a 49% improvement in the Jsc was observed after the DLAR coating deposition. The efficiency of the cells remained low (12 − 14%) however due to a 5% (absolute) reduction in FF and a Voc that remained < 600 mV. Several batches of SP solar cells were produced with both coatings, and the SiNx DLAR coated samples typically exhibited an efficiency advantage of 0.5% absolute. This increased efficiency is not purely due to improved optics, as there are also the electrical benefits of bulk passivation from the hydrogen in the SiNx.

Doshi et al. designed a DLAR coating using two SiNx layers, nAR1 =2.03 and nAR2 =2.42, for use under glass. The DLAR coating performed slightly better than an optimised SiNx 2 2 SLAR coating, with a photocurrent loss of Jpcl =1.95 mA/cm (compared to 2.24 mA/cm for the SLAR coating). Although reflectance losses were low Rw =2.96%, absorption losses arising from the bottom layer (Aw =1.74%) were significant. Doshi et al. claim that even the 0.3 mA/cm2 gain resulting from the DLAR coating is cost-effective, as the wafers can remain in the PECVD chamber for both depositions.339 An additional benefit of such a coating may be improved surface and bulk passivation resulting from the high hydrogen concentration in the bottom layer. 7.3 Varying the Optical Properties of TiO2 167

Figure 7.6: Refractive index and extinction coefficient comparison be- tween the TiO2 SLAR coating used by Solarex and the bottom layer of a 340 SiNx DLAR coating (adapted from Winderbaum et al.).

Graded Index AR Coatings

Johnson et al. postulated that a continually varying refractive index, from 1.5 at the air/film interface to 2.6 at the film/silicon interface, should produce a coating with an average re- flectance of about 3% over most of the visible and infrared wavelengths.349 As a first approxi- mation, four SiNx layers were deposited with the following refractive indices and thicknesses: nAR1 =1.83, dAR1 = 23 nm, nAR2 =1.92, dAR2 = 31 nm, nAR3 =2.20, dAR3 = 24 nm, and 2 nAR4 =3.10, dAR1 = 19 nm. This SiNx MLAR coating was used in the fabrication of a 4 cm 2 solar cell with a Jsc of 34.8 mA/cm .

7.3 Varying the Optical Properties of TiO2

The optical constants of TiO2 films deposited in this work were determined using spectro- scopic ellipsometry (SE), an optical characterisation technique introduced in Section 4.6. A distinct advantage of SE is that the n and k of single- or multi-layer films can be determined relatively easily over a wide range of wavelengths, in our case 350 − 1150 nm. A number of other researchers have used SE to determine the optical properties of TiO2 thin films, as well as other information, such as the void content of the TiO2 thin films and the distribution of the void content across the film.82, 128, 140, 144, 150, 161, 353 Results from these works were dis- cussed in Sections 2.3.3 (optical properties), and Section 2.3.6 (void content). Results from this work will be presented, indicating the variation in optical properties of TiO2 thin films achieved by altering the deposition and annealing conditions (temperature and ambient). 168 7. TiO2 Antireflection Coatings

7.3.1 Deposition Temperature

Experiment

Before designing TiO2 AR coatings, it was necessary to determine the range of n and k values that could be achieved with the CVD system. The aim of this experiment was to ◦ ◦ deposit a range of TiO2 films from the minimum (150 C) to maximum (450 C) deposition temperature. The substrates used in this work were polished 0.1 − 10 Ω cm n-type FZ silicon wafers. TiO2 depositions were performed using a CVD system, and the deposition time and relative humidity (RH) for each sample ranged from 8 − 14 min and 9 − 15.4%, respectively. The higher RH values were for the three samples deposited at ≤ 250◦C. Two additional samples were deposited at 250◦C and 450◦C at a later time in order to determine the repeatability of results. During these depositions the RH was lower at 7 − 8% and film depositions took 21 − 23 min.

SE measurements were performed on each sample in the wavelength range 350 − 1150 nm and at 65 − 80◦ in 10 nm and 5◦ intervals, respectively. The beam size on the sample was about 3 mm in diameter. Sample scans took about 2.5 hr each. The ellipsometric data Ψ and ∆ was modelled using a Lorentz double-oscillator (see Equation 4.7 in Section 4.6.3) using version 3.361 of the software package WVASE32.62

Results and Discussion

Figure 7.7 shows the sample structure used to model the TiO2 films on silicon. The optical constants for the silicon substrate were taken from Green.8 An effective medium approxima- tion (EMA) layer, discussed in Section 4.6.4, comprised of 50% void and 50% TiO2 was used to model the surface roughness in the films. The optical properties of the TiO2 fraction of the EMA layer were coupled to the optical constants of the dense TiO2 layer below it. The thickness of the silicon wafer was fixed at 300 µm.

The fitting parameters used to model the dielectric constant function of this film are shown in Table A.1 (see Appendix A). The modelling parameters for all TiO2 films discussed in this section can also be found in Appendix A. Figures 7.8(a) and (b) demonstrate the ex- cellent agreement reached between the measured and modelled values of the ellipsometric parameters Ψ and ∆. The quality of the fit indicates that it is not necessary to model an inhomogeneity (gradient) in refractive index of the TiO2 layer, as has been required by other researchers.128, 160, 353 Notably, this sample (TO-2-2) exhibited the poorest mean-squared- error (MSE) out of all the samples in this experiment, however the fit to the experimental data is still excellent. In the regions where cos(∆) ≈ 1 it is known that the extinction coef- ficient k has a value close to zero.128 This agrees with the results of many other researchers that TiO2 is essentially transparent at wavelengths greater than 400 nm (see Section 2.3.3). 7.3 Varying the Optical Properties of TiO2 169

air

surface roughness layer dsurf (EMA: 50% void / 50%TiO2)

TiO2 dTiO2

silicon dSi

air

Figure 7.7: Schematic diagram of the film structure used for analysing the SE data using WVASE32.62

The variation of the TiO2 refractive index (n at λ = 600 nm) for a range of deposition tem- peratures is shown in Figure 7.9(a). As expected, n increases with Tdep due to densification and crystallisation processes. The exact values for n, as well as the other parameters plotted in Figure 7.9(b), extinction coefficient (k at λ = 400 nm) and the ratio of the surface layer dsurf thickness to TiO2 layer thickness ( ), are included in Table A.2 in Appendix A. Since the dTiO2 optical constants of both the surface roughness and dense TiO2 layers are coupled, trends observed in n and k for the TiO2 layer are applicable to both layers. However, it should be noted that the absolute n and k values for the surface roughness layer will greatly reduced due to the 50% void incorporation. There is no clear trend in the extinction coefficient (see Figure 7.9(b)) with k values spanning an order of magnitude, and sample TO-2-2 exhibiting a much higher k than other films.

The quantity dsurf , a result of the optical modelling, is used as a parameter to tie together the dTiO2 physical roughness and the optical properties of the modelled layers. Although, SE analysis provides a unique solution for the optical properties and layer thicknesses, the chosen EMA void fraction will also influence the reulting layer thicknesses. Therefore, while the the modelled thicknesses may not exactly match the measured layer thickness, the normalised parameter dsurf serves as a useful indicator as to the behaviour of the surface roughness dTiO2 layer.

◦ The dramatic increase in dsurf for T ≥ 300 C corresponds with the transformation of amor- phous TiO2 to the polycrystalline phase of anatase. Polycrystalline thin films are known to exhibit much greater surface roughness than amorphous films.82, 138 The reason for the reduction in surface roughness from the sample deposited at 200◦C to that of the film de- 170 7. TiO2 Antireflection Coatings

Figure 7.8: Measured (symbols) and modelled (lines) values for (a) cos(∆) and (b) tan(Ψ) for sample TO-2-2.

posited at 300◦C is not clear. It is possibly related to a more complete reaction of the TPT ◦ vapour, as it was observed that the films deposited at Tdep ≤ 250 C were soft and could be easily scratched with plastic tweezers.

The dispersive relations for the refractive index and extinction coefficient of the dense TiO2 layer are illustrated in Figures 7.10(a) and (b). The trends in n and k for the surface roughness layer are identical as the optical constants are coupled to the dense TiO2 layer

(these are included for each TiO2 film in Appendix A). The refractive index curves in Figures 7.10(a) are nearly flat at longer wavelengths, and n rapidly increases at wavelengths less that 7.3 Varying the Optical Properties of TiO2 171

Figure 7.9: Measured parameters for the samples deposited at 150 − 450◦C: (a) Refractive indices and (b) extinction coefficients of the surface roughness layer and the dense TiO2 layer, and the ratio of these layer thicknesses. The lines between symbols are provided as a guide for the eye only.

500 nm due to the approaching bandgap. Only sample TO-2-2 exhibits a dispersive curve with a lower steepness than other samples. The reason for this is not known, but may be due to variations in the amorphous phase of TiO2. As mentioned earlier, there is no clear trend for the extinction coefficient values plotted in Figure 7.10(b) and Figure A.2(b). Sample TO-2-2 deposited at 250◦C exhibits the lowest k, while sample TO-2-4 (300◦C) possesses the highest k. Other researchers have also observed that there is not a direct relationship 172 7. TiO2 Antireflection Coatings

148 between k and Tdep, as presented in Section 2.3.3.

Figure 7.10: Dispersive relations for the (a) refractive index and (b) extinction coefficient of the TiO2 layer of samples deposited using CVD at a range of deposition temperatures.

It is possible that the samples deposited at lower temperatures have retained some organic (OH) groups from the organometallic precursor. It was noted during the depositions that the relative humidity was significantly higher for these samples. It is known that temperatures ◦ 46 of 450 C are required to drive off all OH-groups from within the TiO2 film. However, it is not known exactly how these incorporated OH-groups would explain the variations in k observed in the TiO2 films. 7.3 Varying the Optical Properties of TiO2 173

Two additional samples were deposited at later date in order to check the repeatability of n and k values from the TiO2 films. The dispersive refractive indices and extinction coefficients for the TiO2 layer are plotted in Figure 7.11. Large differences in the optical properties can be seen between the older and newer samples. At a wavelength of 600 nm, the refractive index is 0.059 (or 3.1%) higher at 250◦C and 0.128 (or 5.7%) greater at 450◦C.

Figure 7.11: Dispersive relations for the refractive index (main graph) and extinction coefficient (inset graph) of the TiO2 layer of samples de- posited using CVD at 250◦C and 450◦C.

The most striking variation in the extinction coefficient is that the samples deposited at 450◦C possess the lowest (TO-2-8) and highest (TO-3-1a) k values. It is believed that the much slower deposition rate of samples TO-3-8a and TO-3-1a is responsible for the higher n values observed. The longer time spent at the deposition temperature will enable the

TiO2 grains to sinter and densify somewhat, increasing the refractive index. This hypothesis does not agree, however, with the rule-of-thumb that a slowly deposited film leads to high transparency and low n, whereas high deposition rates result in a high n but also high k.277 The fractionally lower relative humidity present during the later depositions is not likely to have had such an impact on the refractive indices of the films. Another possibility is that the TPT vapour was impinging at a slightly shallower angle in the second round of depositions. From past experience, it has been noted that a slight variation in “deposition angle” can drastically alter the deposition efficiency (the fraction of TPT vapour that is incident upon the substrate), and therefore the deposition rate. It can be imagined that a vapour-stream 174 7. TiO2 Antireflection Coatings impinging on the substrate at a shallower angle could result in less cooling occurring, thus bringing about an effective rise in the deposition temperature. The latter explanation is the most likely mechanism responsible for the large increase in n observed.

7.3.2 Annealing Temperature

Experiment

◦ For this experiment, all TiO2 samples were deposited at 450 C. The relative humidity re- mained between 7−11% for all samples, while the deposition times ranged from 14−17 min. The wafers were each cut into three pieces using a laser. One piece of each sample was retained, while the other pieces were annealed in a quartz tube furnace at 450 − 1050◦Cin a nitrogen (N2) ambient for a period of either 1 hr or 6 hr.

Results and Discussion

The variation of the TiO2 refractive index (n at λ = 600 nm) for a range of annealing temperatures is shown in Figure 7.12(a). The refractive index is roughly constant up until ◦ ◦ 700 C for the films annealed at both 1 hr and 6 hr. At Tann ≥ 800 C, there is a clear trend of increasing refractive index with increasing annealing temperature. This increase in n at higher temperatures is caused by an increasing rutile content in the TiO2 films. The films annealed at 450−700◦C exhibit very similar n values, and it is anticipated that these samples are anatase only. The phase transformation between anatase and rutile can be seen at 800◦C. This sample (TO-4-4c) is thus a mixture of anatase and rutile phases. The samples annealed at 900− 1050◦C possess significantly higher refractive indices, due to a greater rutile content in these films. The samples annealed at 800 − 950◦C for 6 hr exhibit the largest increases in n, 0.135, 0.066 and 0.069 greater than the equivalent samples annealed for 1 hr, respectively. Therefore, it can be concluded that at these temperatures, the crystallisation process is limited by time. The fact that the refractive indices of these samples did not reach the 2.6 ◦ margin of the samples annealed at Tann ≥ 1000 C suggest that the phase transformation is not complete and that these films remain an anatase/rutile mixture.

The refractive indices of the surface roughness layers, plotted in Figure 7.12, are in the range 1.501 − 1.747. These refractive indices are similar to the range represented by SiO2 and Al2O3, both of which are often used as the top layer in a DLAR coating. This suggests the possibility that a single TiO2 deposition step could be used to form a “pseudo”-DLAR coating. The term pseudo-DLAR coating is used as it is not possible to independently optimise the thickness or refractive index of the surface roughness layer - its properties are inherently linked to the denser TiO2 layer. However, the performance of a pseudo-DLAR coating may offer a performance enhancement over a SLAR coating while avoiding the more 7.3 Varying the Optical Properties of TiO2 175

Figure 7.12: Measured parameters for samples annealed at 450 − 1050◦C for a period of 1 hr and 6 hr: (a) refractive indices, and (b) extinction coefficients (only 1 hr data shown) of the surface roughness layer and the dense TiO2 layer, and the ratio of these layer thicknesses. The lines between symbols are provided as a guide for the eye only.

complex processing requirements of a regular DLAR coating. This idea will be explored in greater depth in Section 7.4.

There is a general trend of the extinction coefficients of the films annealed for 1 hr increasing at annealing temperatures greater than 500◦C, as displayed in Figure 7.12(b). The exact explanation for the samples annealed at 900◦C and 950◦C exhibiting the highest k values remains unclear. This trend was observed with samples annealed for 6 hr as well (see Ap- 176 7. TiO2 Antireflection Coatings pendix A). Therefore, as the samples annealed at the same temperature are pieces of the same deposited film the variation in k is attributed to slight fluctuations in the deposition conditions, which results in the films sintering and crystallising slightly differently. It is not expected to be due a result of the SE analysis, due to the use of a Kramers-Kronig consistent optical model and the excellent MSE values achieved, as shown in Table A.5. Hovel also found that a general trend of increasing k with increasing deposition temperatures existed in spray-deposited films, however the occasional film did not fit this trend (see Figure 2.11).148 ◦ It is interesting to note that while films annealed at Tdep < 700 C have very similar n values, their extinction coefficients vary significantly. It is possible that this disparity between n and k is due to the existence of a small rutile fraction in the films or increased disorder in the films with the impending phase transformation.

dsurf ◦ The thickness ratio remains high for Tann ≤ 800 C, as shown in Figure 7.12(b), and is dTiO2 similar to the values observed for non-annealed films deposited at 450◦ (refer Figure 7.9(a)). ◦ At annealing temperatures greater than 900 C, both the surface layer thickness dsurf and dsurf decrease rapidly. The reduced surface roughness at higher annealing temperatures dTiO2 is attributed to sintering of the TiO2 crystallites. The samples annealed for 6 hr exhibit dsurf ◦ slightly reduced values, and reach a minimum value at Tann ≥ 1000 C. SEM images dTiO2 (refer Figure 7.13, later in this section provide independent confirmation of decreased surface roughness for annealed films.

An increase in n and k, and the reduction in the thickness ratio dsurf is expected due to dTiO2 two mechanisms, crystallisation and sintering. The effect of sintering is demonstrated in the scanning electron microscopy images in Figure 4.17. These are higher resolution images of the same samples discussed in Section 4.8.3. After annealing the as-deposited (a) films for 1 hr (b), the crystallites have become rounded and the grains are observed to agglomerate. Annealing for 6 hr results in the majority of grains sintering together, although the largest voids in between the grains remain “unbridged”. After a total of 22 hr annealing at 1000◦C an almost continuous film is formed, with only small voids, typically 40 nm in diameter, remaining. The sintering results in significant densification of the TiO2 layer, with the film thickness decreasing from 79 nm in (a) to 64 nm after step (d). The increase in film density, brought about by the reduced film thickness, is responsible for a small increase in ◦ the TiO2 refractive index. The grain sizes observed in the films deposited at 450 Care roughly 30 nm across, in agreement with nanocrystalline anatase thin films deposited by other researchers.354, 355

An additional mechanism taking place during annealing is crystallisation. Crystallisation is primarily responsible for the large increase in n and k presented in Table A.6. At annealing temperatures less than 800◦C, no significant increase in n is observed. Annealing at 450 − ◦ 700 C will transform any amorphous TiO2 fraction into the anatase phase. However, the films are already anatase as-deposited, and the stable n values indicate that no significant ◦ amorphous fractions were present prior to annealing. For Tann ≥ 800 C, a significant increase 7.3 Varying the Optical Properties of TiO2 177

Figure 7.13: High resolution SEM images of TiO2 films deposited by CVD at 450◦C: (a) as-deposited; (b) 1 hr anneal (load 800◦C and ramp to 950◦C); (c) 6 hr anneal at 1000◦C; and (d) 22 hr anneal at 1000◦C.

in n is noted, reaching a maximum of about 2.6 at a wavelength of 600 nm. This increase is due to the transformation of the anatase phase to the rutile phase. This transformation temperature is in agreement with other works, reviewed in Section 2.2.1. A trend in the k values, determined at λ = 400 nm, is not so clear from Table A.6.

Keddie et al. noted that the higher TiO2 film densities were achieved with increased heating rate.356 The higher film density was attributed to delaying the crystallisation, which allowed 178 7. TiO2 Antireflection Coatings greater film densification to occur prior to crystallisation. It was also emphasised that pre- mature crystallisation can lead to films that are only partially sintered, because diffusive sintering mechanisms are much slower than viscous sintering mechanisms.356 This implies that the densest films may possibly be achieved by depositing amorphous TiO2 films, followed by annealing (sintering) the films at a temperature less than 300◦C to avoid crystallisation into the anatase phase, and finally annealing (crystallising) the samples at a higher tem- perature. Fitzgibbons et al. commented that the crucial parameter determining the final

TiO2 film structure was simply the maximum processing temperature (either deposition or annealing temperature).67 It was also noted that rutile could be achieved directly by per- forming the deposition at moderate temperatures. This is in agreement with the observation of other researchers that the processing temperatures required to convert an anatase film into a rutile film are significantly higher than the temperature required to deposit a rutile film directly.73–75

7.3.3 Deposition Ambient

Experiment

The aim of this experiment was to determine whether the presence of water vapour (H2O) would alter the optical properties of the TiO2 films. For this experiment, TiO2 samples were ◦ ◦ deposited at 250 C and 450 C, both with and without H2O. The water vapour was supplied by heating DI water, contained in a quartz bubbler, to 100◦C. The relative humidity during depositions was about 30%. Certain pieces of the samples were annealed in N2 at either 450◦Cor1000◦C for a period of 1 − 6 hr. One sample was loaded in the annealing furnace in oxygen (O2) for the first 5 min, before the gas flow was changed to N2 for the next 55 min. The optical properties of all of the films in this section were measured using SE, and the results are given below. The fitting parameters used to model the samples are provided in Table A.7 in Appendix A.

Results and Discussion

The refractive indices of the films deposited in the presence of water vapour were typically low at n =1.726 − 2.194. The samples deposited at 250◦C exhibit the lowest refractive indices, regardless of annealing time and temperature. In fact, the film annealed at 1000◦C for 6 hr has a very similar refractive index to the film annealed 450◦C for 1 hr. However, Table 7.2 indicates that the former film has an extinction coefficient many orders of magnitude greater than the latter film. This increase in k can partly be attributed to the transformation of the film from anatase to rutile. It is also interesting that the sample TO-4-1c, deposited at 450◦C, exhibits a significantly higher n and k than the film annealed at 450◦C. This may indicate that the as-deposited amorphous phase is more porous than the as-deposited 7.3 Varying the Optical Properties of TiO2 179 anatase phase, and that the porosity can not be reduced during subsequent anneals. This observation is in agreement with the results of other researchers. Wong et al. also reported ◦ that TiO2 films deposited at 265 C were porous and hence the rearrangement of the atoms by a sintering process is limited.130 This phenomenon is also supported by the marked increase in n observed for the films from this work that were deposited at 450◦C and subsequently annealed, as shown in Figure A.7(a).

Table 7.2: Thickness, refractive index (at λ = 600 nm) and extinction co- efficient (at λ = 400 nm) of surface roughness and TiO2 layers deposited ◦ ◦ using CVD at 250 C and 450 C with H2O vapour. Some samples were ◦ ◦ annealed at either 450 Cor1000 CinN2 for a period of 1 hr or 6 hr.

Surface Layer TiO2 Layer T T d nkd nk Sample dep ann surf TiO2 dsurf Name (◦C) (◦C) (nm) (nm) dTiO2 No Anneal TO-3-6a 250 − 63.7 1.395 0.00037 52.1 1.829 0.00082 1.223 TO-4-1c 450 − 73.2 1.501 0.00639 32.0 2.004 0.00733 2.288

1 hr Anneal in N2 TO-3-6b 250 450 89.6 1.347 0.00022 68.5 1.726 0.00047 1.308 TO-3-5a 450 1000 43.2 1.562 0.03425 74.0 2.194 0.07557 0.584

6 hr Anneal in N2 TO-3-7c 250 1000 68.7 1.350 0.01951 60.7 1.731 0.04231 1.132 TO-3-5b 450 1000† 49.7 1.524 0.06151 80.0 2.111 0.13568 0.621

† Sample TO-3-5b was annealed for 5 min in O2 andsubsequently5hr55mininN2 in the same high-temperature step. This resulted in the growth of 11.1nmofSiO2 at the TiO2:Si interface.

One surprising result is that the sample deposited at 250◦C (TO-3-6a) had a refractive index of 1.829, however after annealing at 450◦C (TO-3-6b) this decreased to n =1.726.

The samples deposited at 250◦C (TO-3-6a, TO-3-6b an TO-3-7c) exhibited very low re- fractive indices. This reduction in n is attributed to a large volume of water vapour being contained within the film during deposition. The water vapour is then subsequently driven- off by heating of the sample. The porosity of sample To-3-6b, determined using Equation 2.3 and using nb =2.532 (dense anatase), is 63.4%. As mentioned in Sections 2.3.6 and 2.3.4, highly porous films can absorb water vapour and this is known to increase the refractive index of porous samples.155, 161, 163, 164 The voids in the film become essentially “filled” with n . water ( H2O =133) and the measured refractive index can be significantly higher than if the film was measured in a vacuum. The possibility of this effect could not be precluded in the SE measurements, and the samples were measured in air.

The surface roughness layer of the samples deposited in the presence of water vapour was 180 7. TiO2 Antireflection Coatings

often thicker than the dense TiO2 layer, as shown in Table 7.2. The films with the highest n values are both pieces of the same sample (TO-3-5), and these films also exhibit the lowest dsurf values. This is in agreement with the films possessing a higher density than the others. dTiO2 In general, the large dsurf values and the low refractive indices of the TiO2 layer indicate that the films are quite rough and porous.

Figure 7.14(a) is an SEM image of sample TO-3-6a, deposited with water vapour at 250◦C. Larger clusters of about 60 − 70 nm across are comprised of smaller grains with a diameter of roughly 10 nm. The large voids in between the clusters are also immediately apparent. Lottiaux et al. reported similar observations, noting that larger “blocks” of about 40 nm in size were comprised of smaller blocks that were about ten times smaller in size.242 The formation of the smaller blocks was found to be in agreement with nucleation theory and the growth of amorphous thin films. Additionally, SEM images showed that these blocks do not change with deposition conditions, and that deep cracks separate the large blocks.

◦ Figure 7.14: SEM images of (a) TiO2 films deposited at 250 C (TO-3-6a) and (b) subsequently annealed at 1000◦C 6 hr (TO-3-7c).

Figure 7.14(b) shows the same film after annealing for 6 hr at 1000◦C. The clusters and grains have undergone significant densification, however the large voids are still present. This indicates that sintering has only occurred when there was TiO2 material directly adjoining the grains. The refractive index of the sample annealed for 6 hr at 1000◦C (TO-3-7c) was very similar to the sample annealed for 1 hr at 450◦C (TO-3-6b). Assuming that after a 6 hr anneal the film is converted to 100% rutile, this demonstrates that films can be deposited that exhibit certain properties of rutile (such as excellent chemical resistance) while maintaining a low refractive index. 7.3 Varying the Optical Properties of TiO2 181

7.3.4 Annealing Ambient

Experiment

The aim of this experiment was to determine whether the gas ambient during annealing would later the optical properties of the TiO2 films. Other researchers have reported that the optical properties of TiO2 can vary depending on the whether the gas flowing through 150 ◦ the furnace is O2 or N2. Therefore, samples were deposited at 450 C and annealed in a quartz tube furnace at 1000◦C for a total time of 1 hr or 6 hr. The only variable was whether the samples received 5 min O2 at the start of the furnace step or not. The relative humidity remained between 7 − 11% for all samples. In all cases the total time the sample spent in the furnace is either 1 hr or 6 hr.

It was demonstrated in Section 5.3.1 that a thin SiO2 could be grown at the TiO2:Si inter- face by performing a brief oxidation (dry O2). Therefore, the final section of work was to determine how the optical properties of the TiO2 film changed with increase oxidation times and temperatures. Five samples were deposited at 450◦C(15− 23 min deposition time and ◦ RH ≈ 7%) and loaded individually into a quartz tube furnace at 800 CinanO2 (3.2 slpm) ambient. Once loaded, the furnace was set to ramp to 1000◦C, and the wafers were removed after 1 min, 2 min, 4 min, 8 min and 16 min.

The optical properties of all of the films in this section were measured using SE, and the results are given below. The fitting parameters used to model the samples in this section are provided in Tables A.8 and A.9 in Appendix A.

Results and Discussion

To compare the effect on the annealing ambient, pairs of samples were annealed either in pure N2 for 1 hr or 6 hr, or initially 5 min O2 followed by N2 for a total processing time of 1 hr or 6 hr. Table 7.3 details the thickness and optical properties of the various layers.

It was necessary to include an SiO2 layer at the TiO2:Si interface in the model to achieve d the best fit to the experimental data. The thickness of this layer SiO2 varies somewhat, from 6.2 − 10.9 nm. Slight variations in the SiO2 thickness may be caused by TiO2 film inhomogeneities. From Table 7.3 it can be seen that one pair of samples (TO-3-3b and TO-3-3c) possess virtually identical refractive indices at 600 nm. However, for the remaining pairs (TO-3-5a and TO-3-5b, and TO-4-7c and TO-3-3a), the films annealed only in N2 exhibited higher refractive indices. The similarities between the n values of the former pair and the difference in the latter pair is clearly illustrated in Figure A.8(a).

Figure A.8(b) indicates that the extinction coefficient values for all samples are fairly similar, except the sample TO-3-3a, which is significantly higher than the others. Kamataki et al. 150 observed slightly lower n and k values for films annealed in O2. This research group 182 7. TiO2 Antireflection Coatings

Table 7.3: Thickness, refractive index (at λ = 600 nm) and extinction co- efficient (at λ = 400 nm) of surface roughness and TiO2 layers deposited using CVD at 450◦C. Some samples received a 5 min oxidation prior to ◦ all samples being annealed at 1000 CinN2 for a total period of 1 hr or 6 hr.

Surface Layer TiO2 Layer d d nkd nk Sample 5 min SiO2 surf TiO2 dsurf dTiO2 Name O2? (nm) (nm) (nm) 1 hr Anneal TO-4-7c ✗ − 14.2 1.722 0.04047 55.0 2.547 0.08962 0.258 TO-3-3a ✓ 6.2 30.5 1.697 0.05312 49.7 2.491 0.11766 0.614 TO-3-5a† ✗ − 43.2 1.562 0.03425 74.0 2.194 0.07557 0.584 TO-3-5b† ✓ 10.9 49.9 1.524 0.06130 80.0 2.111 0.13517 0.624 6 hr Anneal TO-3-3b ✗ − 10.5 1.761 0.02513 64.0 2.633 0.05565 0.164 TO-3-3c ✓ 10.7 22.0 1.761 0.04309 46.9 2.632 0.09540 0.469

† Samples TO-3-5a and TO-3-5b were deposited in the presence of H2O vapour.

attributed this to the growth of SiO2 at the TiO2:Si interface of O2 annealed samples. The ◦ analysis of Kamataki et al. determined that after 1 hr at 500 Can11.6 nm-thick SiO2 layer had grown at the interface. This SiO2 thickness is similar to the samples in this thesis, d however the annealing temperature is halved. It would seem unlikely that a similar SiO2 would grow at a temperature of 500◦C. It may be possible that a significant native oxide existed prior to TiO2 film deposition instead. Examination of the results of Kamataki et al. in Figure 2.17 indicate that an increased refractive index is primarily observed at wavelengths less than 500 nm. In conjunction with the increased k at shorter wavelengths, it would seem to suggest that the film annealed in N2 may possess a small rutile fraction. The results from this thesis also exhibit increased n values for films annealed for 1 hr. After 6 hr annealing, refractive indices were identical. Therefore, it postulated that oxygen may play a role in slightly retarding the transformation from anatase to rutile.

The following results are for the samples deposited at 450◦C and loaded into an oxidation furnace at 800◦C. The furnace temperature was set to ramp to 1000◦C. The samples received separate oxidations of varying length - 1, 2, 4, 8 and 16 min - and were unloaded without ramping the furnace temperature down. Table A.10 (in Appendix A) and Figure 7.15 sum- marise the measured parameters, including refractive index, extinction coefficient, SiO2 layer thickness and dsurf . dTiO2 The refractive indices of the samples are observed to increase with increased oxidation time. From the n values, it is believe that the transformation from anatase to rutile begins with the sample oxidised for 8 min. There is also a large increase in k for samples for times 7.4 Development of Novel TiO2 AR Coatings 183 longer then 8 min, as shown in Figure A.9(b) (refer Appendix A). It remains unclear as to why the thickness ratio dsurf steadily increases with increased oxidation time (and therefore dTiO2 temperature). It is postulated that crystallisation is occurring first and that densification has not had a chance to proceed yet. This is in agreement with the observations of Keddie et al., who noted the rate of crystallisation is much greater than the rate of densification, and that a high density film (corresponding to a thin surface roughness layer) is achieved by delaying crystallisation.356

7.4 Development of Novel TiO2 AR Coatings

The achievement of a wide range of refractive indices by altering the deposition and annealing conditions suggested the possibility of creating several novel types of AR coatings using only

TiO2. This idea was supported by the fact that a single deposition resulted in a denser

TiO2 layer (n =1.726 − 2.633 at λ = 600 nm) underneath a surface roughness layer (with n =1.347 − 1.761 at λ = 600 nm). Although the thickness of the surface roughness layer could not be independently, the behaviour of dsurf as a function of deposition temperature was known. This section presents new designs of SLAR and DLAR coatings based on these

CVD-deposited TiO2 thin films.

7.4.1 Single-layer TiO2 AR Coatings

Measurements on the above TiO2 samples were performed in order to determine whether the surface roughness layer would act as a “pseudo-DLAR coating”. A pseudo-DLAR coating is defined as an optical coating that can reducing the amount of reflected light by more than is possible with an SLAR coating, but is still the result of a single film deposition. It could also be regarded as the first step from a SLAR coating towards a graded refractive index.

Experiment

The deposition of the samples used in these measurements was previously described in the ‘Experiment’ section of Sections 7.3.1, 7.3.2, 7.3.3 and 7.3.4. Reflectance measurements were performed with a Varian Cary 5G spectrophotometer in the wavelength range 210−1200 nm. A lamp was used as the light source for wavelengths less than 350 nm, while a tungsten-halogen lamp was used at longer wavelengths.

The data was modelled using the Lorentz oscillator model presented in Section 4.6.3. The pseudo-DLAR coating model included a surface roughness layer and a dense TiO2 layer, while the single layer model consisted of one dense TiO2 layer. 184 7. TiO2 Antireflection Coatings

2.4 (a)

2.2 Surface roughness layer =600nm) λ Dense TiO layer 2.0 2 (at (at n 1.7

1.6

Refractive Index, Refractive Index, 1.5 0246810121416 Oxidation Time, t (min) ox

0.56 10 0.54 (b)

2

TiO (nm) 8 2

d 0.52

/ SiO

d

surf surf d 0.50 6

0.48 4

0.46 2 Layer Thickness, Layer 2 Thickness Ratio, 0.44 SiO 0 0 2 4 6 8 10 12 14 16 Oidti Ti t (i) Figure 7.15: Measured parameters for samples oxidised for 1 − 16 min: (a) refractive indices, and (b) thickness ratio of the surface roughness layer and dense TiO2 layer and the SiO2 layer thickness grown. The lines between symbols are provided as a guide for the eye only. 7.4 Development of Novel TiO2 AR Coatings 185

Results and Discussion

The reflectance spectra of a typical TiO2 film, sample TO-3-1b, is shown in Figure 7.16. The effect of the surface roughness layer can be seen, broadening the reflectance minimum and significantly reducing the short wavelength reflectance. The weighted average reflectance for the pseudo-DLAR coating is Rw =8.95% over the wavelength range 350 − 1150 nm, compared to Rw =10.54% for the best single-layer fit. Thus, the coating is behaving like n2 n n2 n a DLAR coating with the refractive indices satisfying the relationship AR1 Si = AR2 0.

The reduction in Rw can be attributed solely to the presence of the surface roughness layer with an intermediate refractive index, since the extinction coefficient of both the dense TiO2 layers was virtually identical.

Figure 7.16: Reflectance of sample TO-3-1b in air, modelled with a single layer (dotted line), and EMA/dense layer (solid line).

It should be noted that, even though sample TO-3-1b exhibited one of the lowest reflectances, its refractive index was greater than desirable (nAR1 =2.08) for a AR coating in air. However, from Table A.2 it can be seen that the thickness of the surface roughness layer decreases significantly for deposition temperatures lower than 450◦C. This highlights the limitation of the pseudo-DLAR coating - that the thickness of the surface roughness layer cannot be independently controlled.

Figure 7.17 models the performance of sample TO-4-5c under 2 mm of B270 Crown glass and 1 mm ethyl-vinyl-acetate (EVA) encapsulant. The optical models for the B270 Crown glass and the EVA layer were taken from Nagel et al.45 Although, the extinction coefficient of the EVA very low (k =7.7 × 10−6 at 400 nm), the large thickness results is the majority of the light at wavelengths less than 400 nm being absorbed in the layer. Again, the optical performance is seen to be enhanced slightly by the presence of the TiO2 surface roughness layer, primarily at short wavelengths. This reduces the weighted average reflectance from

Rw =8.95% to Rw =8.42%. In addition, the inclusion of a TiO2 film with a refractive index 186 7. TiO2 Antireflection Coatings

of about 2.4 results in a 0.5% (absolute) reduction in Rw compared to the case in air.

Figure 7.17: Reflectance of sample TO-4-5c under glass, modelled with a single layer (dotted line), and EMA/dense layer (solid line).

7.4.2 Double-layer TiO2 AR Coatings

This experiment was aimed at demonstrating that a DLAR coating, comprising two TiO2 layers, could be achieved using the CVD-deposition technique. The existence of a DLAR coating can be confirmed by the observance of two minima in the reflectance spectrum of the coating.

The refractive indices of air, glass and silicon are 1.0, 1.52 and 3.939 at a wavelength of 600 nm. Using Equations 7.8 and 7.6, the optimum refractive indices and thicknesses for a DLAR coating can be determined for a silicon solar cell in air and under glass at a design wavelength of λ0 = 600 nm. In air, the optimum refractive indices (and thicknesses in parentheses) of the two layers are nAR1 =1.58 (94.9 nm) and nAR2 =2.49 (60.2 nm), while under glass these values increase to nAR1 =2.07 (72.5 nm) and nAR2 =2.86 (52.4 nm). These refractive indices satisfy the requirements of a DLAR coating in order to achieve zero reflectance at two wavelengths either side of λ0, nAR1 nAR2 = n0 nSi.

As can be seen from Table A.6, sample TO-4-5b exhibited a refractive index of 2.489, very close to that required for the bottom layer of the DLAR coating in air. As this layer was annealed and had undergone densification and crystallisation processes, it was necessary to estimate the degree of thickness reduction. An ellipsometer (λ = 632.8 nm) in the labora- tory was used to perform quick measurements during sample processing. Sample TO-4-5b possessed a thickness and refractive index of 95.1 nm and 1.991, respectively, after depo- sition. After annealing for 6 hr at 900◦C, these values changed to 57.3 nm and 2.509 (at

λ = 632.8 nm). To achieve a minimum at λ0 = 600 nm with this film (nAR =1.991), a thick- ness of 75.2 nm is required. The thickness of TO-4-5b after annealing is very close to what 7.4 Development of Novel TiO2 AR Coatings 187 is required as a final thickness for the annealed bottom layer. Therefore, when depositing the bottom layer of the DLAR coating, this film will appear thicker than typically desired dark blue, possibly exhibiting a light blue or yellowish colour.

The lowest refractive indices for the dense TiO2 films observed in this work (n ≈ 1.73) were achieved by performing depositions at 250◦C in the presence of water vapour, followed by a subsequent annealing step of 450 − 1000◦C (see samples TO-3-6b and TO-3-7c in Table 7.2). The refractive indices (and thicknesses) of these films were 1.479 (137.6 nm) and 1.551 (157.6 nm), respectively, when measured with the λ =632.8 nm ellipsometer. This piece of ellipsometer is only able to model a single film only, and cannot model surface roughness layers. Therefore, these deposition settings for the upper layer of the DLAR coating. The deposition of the upper layer would be stopped when a dark blue colour was observed in the

TiO2/TiO2 stack.

Experiment

The wafers used for optical experiments were three 0.1 − 10 Ω cm n-type polished float zone (FZ) Si(100) wafers and one chemically-polished, mc-Si wafer (Eurosil, F43, p-type, 1.5 Ω cm). All wafers were cleaned in RCA1/RCA2/HF before being thoroughly rinsed in ◦ DI water and dried. The bottom TiO2 layer was deposited at 450 C, with deposition times and relative humidities ranging from 12 − 23 min and 7 − 11%, respectively. Film colours ranged from dark blue to yellow. All samples were then cleaned in H2SO4:H2O2:H2O (1:1:5) and RCA2 before being rinsed and dried.

Two of the polished FZ wafers (TO-5-5 and TO-5-6) and the mc-Si wafer (TO-5-3) were ◦ then placed in a quartz tube furnace for annealing at 900 Cfor6hrinanN2 ambient. The remaining sample (TO-5-8) received the same amount of annealing time and at the same temperature, however the first 10 min were in O2 before switching the furnace gas to N2. ◦ Subsequently, the upper TiO2 layer was deposited at 250 C with H2O vapour present. The deposition times ranged from 5 − 20 min and the relative humidity was relatively constant at 33 − 34%. As the layers deposited at low temperatures were quite soft, the samples were then annealed for 2 hr at 700◦C.

The reflectance of all samples was measured in the wavelength range 300 − 1200, while spectroscopic ellipsometry measurements were performed in the wavelength range 350−1150 nm. SE data were collected for each sample at angles of 65 − 80◦ in 5◦ steps. The fitting parameters used to model the samples in this section are provided in Table A.11 in Appendix A. 188 7. TiO2 Antireflection Coatings

Results and Discussion

The excellent agreement between the measured (symbols) and modelled (lines) ellipsometric parameters for sample TO-5-5 is shown in Figure 7.18. As noted previously, in regions where cos(∆) ≈ 1 it is known that the extinction coefficient k is very close to zero.128

Figure 7.18: Measured (symbols) and modelled (lines) values for (a) cos(∆) and (b) tan(Ψ) for sample TO-5-5.

The refractive indices, extinction coefficients and thicknesses of the dense TiO2 and surface roughness layers in the DLAR coatings are plotted in Figure 7.19(a), (b) and (c). When dis- played graphically, the approximation of the four modelled TiO2 layers to a graded refractive index coating can be seen. It was attempted to model the layer as a graded refractive index 7.4 Development of Novel TiO2 AR Coatings 189 coating. The void content of each of the four layers was varied, and the optical properties where linked to a bottom layer that determined the optical properties of the film. This model was not successful, and it is postulated that the upper and lower TiO2 films are composed of different phases of TiO2, anatase and rutile, respectively, and possess different optical properties.

The reflectance of three of the TiO2 DLAR coatings is shown in Figure 7.20. The reflectance spectra of sample TO-5-6d is not plotted as it is virtually identical to sample TO-5-5d. All three samples in Figure 7.20 exhibit the characteristic double minimum of a DLAR coating. The noisy data in the range 800 − 1000 nm is due to poor sensitivity in the spectrophotome- ter’s IR detector, however the modelled data clearly provides an excellent fit. The lowest weighted average reflectances achieved for a TiO2 DLAR coating in air in this work is 6.54% and 8.04% for planar c-Si and mc-Si substrates, respectively (see Figure 7.20). While sig- nificantly better than a TiO2 SLAR coating, the performance of this coating is less than optimal due to the high refractive index of the lower TiO2 layer (nAR2 > 2.6 instead of 2.49 as determined at the start of this section).

The most interesting application for a TiO2 DLAR coating would be for planar mc-Si so- lar cells encapsulated under glass. Modelling was performed, using the software package TFCalc,341 with the above coatings to determine the optimum performance that could be expected for this case. The optical models for the 2 mm-thick B270 Crown glass and 1 mm- thick ethyl-vinyl-acetate (EVA) layers were taken from Nagel et al.45 Although, the extinc- tion coefficient of the EVA very low (k =7.7 × 10−6 at 400 nm), the large thickness results is the majority of the light at wavelengths less than 400 nm being absorbed in the layer.

As previously discussed, the ideal refractive index for the lower DLAR coating layer is nAR2 =2.86. As this is not achievable in practice, the highest refractive index obtained in this study was used, nAR2 =2.63 (this can be achieved by annealing the TiO2 layer at 1000◦C for 6 hr). With the reduced refractive index of the lower DLAR coating, the refractive index of the upper DLAR coating also reduced from 2.07 to nAR1 =1.95. Figure 7.21 plots the reflectance, transmittance, and the absorptance of this DLAR coating. The weighted averages of these parameters are Rw =6.98% (including back reflectance), Tw =91.65% and

Aw =2.25%. The reflectance spectra is extremely flat and lies between 4.7% and 7.7% for all wavelengths in the range 410 − 1040 nm. This is excellent, considering that about 4.3% (absolute) is lost in the first bounce off the front surface of the glass. With the inclusion of a 10 nm-thick SiO2 layer at the interface, the average weighted reflectance and absorptance increase to Rw =7.57% and Aw =2.10%, respectively, while Tw =90.33%.

A known advantage of DLAR coatings is that they are less sensitive to variations in their layer thicknesses than SLAR coatings.342, 343, 349 Figure 7.22 plots the reflectance spectra of the TiO2 DLAR coating with ideal layer thicknesses (the same as appears in Figure 7.21), and with both layers 20% thicker and thinner than ideal. The lower TiO2 layer thickness 190 7. TiO2 Antireflection Coatings

Figure 7.19: Optical constants for four various layers of TiO2 DLAR coatings: (a) sample TO-5-5d, (b) sample TO-5-6d and (c) sample TO- 5-8d (N.B. TO-5-8d has a thin SiO2 layer at the TiO2:Si interface. 7.4 Development of Novel TiO2 AR Coatings 191

Figure 7.20: Measured reflectance of three TiO2 DLAR coatings: (a) sample TO-5-5d, (b) sample TO-5-8d and (c) sample TO-5-3d (mc-Si). Note that sample TO-5-8d has a thin SiO2 layer at the TiO2:Si interface. 192 7. TiO2 Antireflection Coatings

Figure 7.21: Reflectance, transmittance and absorptance spectra for the ideal TiO2/TiO2 DLAR coating. Experimental data was used for the dispersive refractive indices of the TiO2 layers.

is 78.0 ± 15.6, while the upper TiO2 layer has a thickness of 49.4 ± 9.9. For the scenario where both layers are too thin , the weighted average reflectance increases from 6.98% to only 7.10%, while thicker layers lead to a slightly higher value of Rw =7.46%.

Figure 7.22: Variation of reflectance spectra with a 20% increase or re- duction in both TiO2 layer thicknesses.

One disadvantage of DLAR coatings is that they more sensitive to thickness of the SiO2 passivation layer thickness than SLAR coatings.344 The results of modelling performed in this work for SLAR and DLAR TiO2 coatings in air is shown in Figure 7.23. This graph shows that the point where a DLAR coating offers no benefit over an SLAR coating is for

SiO2 layers with a thickness of 15 nm or more. As previously discussed, a 10 nm-thick SiO2 passivation layer has been found to achieve excellent results. Therefore, the efficiency of solar cells incorporating a thin SiO2 layer can be expected to increase upon the deposition 7.5 Performance of TiO2 DLAR-Coated Solar Cells 193

of a TiO2 DLAR coating.

Figure 7.23: Dependence of weighted average reflectance on the SiO2 passivation layer thickness for TiO2 SLAR and DLAR coatings in air.

7.5 Performance of TiO2 DLAR-Coated Solar Cells

To predict the performance of solar cells with these DLAR coatings, the weighted average reflectance curves were used as an input to a PC1D simulation. The intensity of the AM1.5 global spectrum was modified to account for absorption in the glass, EVA and TiO2 layers.

The short-wavelength absorptance in the glass, EVA and TiO2 layers is shown in Figure 7.24. This resulted in an intensity of 98.35 mW/cm2, slightly lower than the AM1.5G value of 100 mW/cm2.

The source of the optical losses in the encapsulated solar cell is broken down in individual layers in Figure 7.25 at two different wavelengths, 350 nm and 400 nm. The main source of absorptance is the 1 mm thick EVA layer, with the lower TiO2 absorbing a much smaller fraction of the light. Minimal absorption is observed in the upper TiO2 thin film, while the level of absorption in the glass is close to zero at these wavelengths.

The parameters for the standard buried-contact solar cell, displayed in Table 7.4, were taken from Honsberg et al.23 The IQE of this solar cell is plotted in Figure 7.26.

On 1 Ω cm p-type material, this modelled device resulted in a short-circuit current den- 2 sity of 37.5 mA/cm without an SiO2 passivation layer present, reducing slightly to Jsc = 2 37.2 mA/cm with the inclusion of 10 nm of SiO2 at the TiO2:Si interface. A short-circuit current density of greater than 37 mA/cm2 is excellent value for a planar encapsulated solar cell. This compares favourably to the best buried-contact solar cell manufactured by BP 2 Solar, which achieved a Jsc of 38.3 mA/cm (measured in air) using a silicon nitride SLAR coating on textured crystalline silicon.15 The emitter dark saturation current density was 194 7. TiO2 Antireflection Coatings

Figure 7.24: Intensity of the AM1.5 global spectrum striking the front surface of the glass and the silicon wafer, indicating absorption in the glass, EVA and TiO2 layers.

Figure 7.25: Percentage of optical losses attributable to various layers at wavelengths of 350 nm and 400 nm.

kept constant for TiO2 DLAR coatings both with and without an SiO2 passivation layer 13 2 present. While it is not realistic to expect a J0e of 1 × 10 A/cm with no SiO2 layer, this value was used to indicate the maximum optical performance of the TiO2 DLAR coating. For the device with a surface passivation layer, am open-circuit voltage of 640.6mV was achieved, resulting in a solar cell efficiency of η =18.6%. The fill-factor of this solar cell was 78.1%, which is believed to be achievable for a BC solar cell fabricated on mc-Si substrates.

The high bulk minority carrier lifetime of 100 µs assumes that gettering and/or hydrogenation of the mc-Si wafers has occurred during processing. Gettering is known to occur during the heavy phosphorus groove diffusion,357 while a forming gas anneal performed after the aluminium alloying process has resulted in bulk minority carrier lifetimes in excess of 80 µs on ungettered 1.5 Ω cm mc-Si wafers.35 For mc-Si material that has received no lifetime 7.5 Performance of TiO2 DLAR-Coated Solar Cells 195

Figure 7.26: IQE of a simulated BC solar cells on τbulk = 100 µs material.

Table 7.4: PC1D modelling parameters used to simulate a single-sided buried-contact solar cell (adopted from Honsberg et al.23).

Parameter PC1D Value Surface texturing None p-type wafer resistivity 1 Ω cm Wafer thickness, dSi 350 µm Bulk minority carrier lifetime, τbulk 100 µs 19 −3 Peak emitter doping, N0 1 × 10 cm Junction depth, xj 1 µm Emitter sheet resistance 150 Ω/2 −13 2 Emitter dark saturation current density, J0e 1 × 10 mA/cm 2 Rear surface recombination velocity, SRVrear 1000 cm/s 2 Series resistance, Rs 1Ωcm 2 Shunt resistance, Rsh 100000 Ω cm

enhancement processes, a bulk minority carrier lifetime of 20 µsismoretypical.Onthis material, a BC solar cell with a TiO2 DLAR coating could be expected to achieve an efficiency 2 of 17.4% (Jsc =35.8 mA/cm and Voc = 625.0 mV).

The DLAR coating results represent a 7% improvement over a typical commercial TiO2 SLAR coating, deposited at 320◦C and annealed briefly at 850◦C to simulate firing of the front screen-printed contacts. When modelled together with the electrical parameters for 2 a BC solar cell (with τbulk =100µs), this SLAR coating resulted in a Jsc of 34.7 mA/cm

(Voc =621.3mV and η =16.8%). The refractive index and extinction coefficient of this commercial SLAR coating were n =2.275 and k = 0, at wavelengths of 600 nm and 400 nm, respectively. 196 7. TiO2 Antireflection Coatings

7.6 Conclusions

There are two main candidates for AR coatings in a PV production environment, TiO2 and

SiNx. TiO2 has excellent optical properties and can be deposited at a low-cost. SiNx films are deposited at a higher cost using vacuum-based deposition equipment and exhibit slightly poorer optical properties (higher absorption for a similar refractive index). However, there can also be electrical benefits in using SiNx films deposited using PECVD, due to the large hydrogen content of these films.

It is commonly noted that DLAR coatings have reduced sensitivity to variations in film thickness.342, 343 Johnson et al. noted that a 10% variation of the layer thicknesses in a multilayer coating resulted in only a 1% variation in the reflectance.349 Jellison and Wood noted that the position of the reflectance minimum in the UV spectrum depends primarily on the thickness of the inner layer of the DLAR coating.344 The same layer thickness also affects the IR reflectivity to a much greater extent than the outer layer thickness. The same authors also showed that the outer layer thickness only influences the absolute value of the

UV reflectance minimum. DLAR coatings are much more sensitive to the SiO2 passivation layer thickness, and therefore it is important that the passivation layer be made as thin as possible.344

Doshi et al. explored the idea of using a DLAR coating made from the same materials, 339 in this case SiNx, for use with encapsulated solar cells. While the absorption due to the bottom, high refractive index SiNx layer was high(Aw =1.74), the reduced reflectance

(Rw =2.96%) lead to an overall increase in optical performance. It is claimed that even a small photocurrent gain of 0.3/,mA/cm2 resulting from using the DLAR coating rather than an optimized SLAR coating is cost-effective, as the wafers can remain in the PECVD chamber for both depositions.339 An additional benefit of such a coating may be improved surface and bulk passivation resulting from the high hydrogen concentration in the bottom layer. When a 10 nm-thick SiO2 passivation layer was included underneath a MgF2/SiNx DLAR coating, it was noted that the bottom layer is reduced in thickness by roughly the same amount, while the optimum refractive index for the same layer increases slightly.

All researchers that have investigated multilayer coatings based on SiNx found that the increased extinction coefficient, especially for films with n ≥ 2.3, has resulted in poor optical performance of the coating.45, 339, 340, 349 Therefore, to achieve a practical DLAR coating to improve the performance of planar mc-Si solar cells an alternative dielectric needs to be found.

The optical properties of TiO2 films could be varied greatly, by altering the deposition and annealing conditions (temperature and ambient). TiO2 films with a range of refractive indices from 1.879 − 2.097 (λ = 600 nm) were achieved by varying the deposition temperature from 150−450◦C. Such a trend is not evident in the extinction coefficient values, however all films 7.6 Conclusions 197 exhibit low absorption with k ≤ 1 × 10−2 at a wavelength of 400 nm. The reproducibility between depositions was within ±10% of refractive index. This value is high, and the large variation is attributed to the shallower deposition angles causing less cooling of the substrate, thereby resulting in a higher effective substrate temperature.

By placing the samples in a furnace at 450 − 1050◦C for a period of 1 hr or 6 hr, refractive indices in the range 2.061 − 2.602 (λ = 600 nm) could be obtained. Again, the extinction coefficients of these films did not increase with the same linear trend and, in fact, the highest k values (at 400 nm) were exhibited by films annealed at 900 − 950◦C. However, as the films were deposited individually, there are slight variations in the initial n and k values, and these variations still remain after annealing. The thickness of the surface roughness layer decreased dramatically for annealing temperatures ≥ 900◦C due to sintering processes. TiO2 films annealed for the first 5 min in oxygen exhibited a 6.2 − 10.9 nm-thick

SiO2 layer at the TiO2:Si interface. The refractive indices of samples annealed for 6 hr were the highest observed in this work, at n =2.633. TiO2 samples that underwent a furnace step in an oxygen ambient showed an increasing SiO2 layer thickness for longer processing times. For films deposited in the presence of water vapour, the refractive indices were low at n =1.726 − 2.194 (λ = 600 nm), even after annealing at 1000◦C. An anatase film with a refractive index of 1.726 is calculated to be extremely (63.4%) porous.

A DLAR coating was designed and fabricated, with both layers being comprised of TiO2.A ◦ denser bottom TiO2 layer was obtained by annealing a TiO2 film deposited at 450 C, while a highly porous upper TiO2 layer was achieved by performing the deposition in the presence of water vapour. A distinctive double minima, characteristic of a DLAR coating, was observed in the reflectance spectrum. The best weighted average reflectance measured was 6.54% for a planar crystalline silicon wafers and 8.04% for a planar multicrystalline silicon wafer (both in air). The layer thicknesses and refractive indices were both greater than ideal, indicating that further improvement is possible in the experimental coatings.

Modelling determined that a TiO2 DLAR coating under glass could achieve a minimum weighted average reflectance of 6.1%, while the inclusion of a 10 nm-thick SiO2 passivation layer at the interface increased the Rw slightly to 6.9%. The performance of the DLAR coating was significantly higher than a SLAR, which exhibited an Rw of 8.3% and 8.1%, with and without the SiO2 layer, respectively.

Optical modelling determined that the majority of the absorptance in the short wavelength region occurred in the EVA layer, rather than in the glass or TiO2 layers. A reduced AM1.5G spectrum and intensity was calculated These values, along with parameters describing the electrical performance of a single-sided buried-contact (BC) solar cell, were used as inputs to a PC1D simulation. The maximum Jsc achievable with a planar BC solar cell implementing 2 a TiO2 DLAR coating under glass was 37.5 mA/cm . When a 10 nm-thick SiO2 passivation 2 layer was included at the TiO2:Si interface, this value decreased slightly to 37.2 mA/cm . The 198 7. TiO2 Antireflection Coatings

modelled efficiency for the latter solar cell was 18.6% (Voc =640.6 mV), which is excellent considering the planar nature of the device. The fill-factor for this solar cell is FF =78.1%, which is achievable for a BC solar cell fabricated on mc-Si substrates. Use of the DLAR coating results in a 7% improvement in Jsc over a typical commercial TiO2 SLAR coating on the same device. Chapter 8

Conclusions

8.1 Summary

Overview and Motivation

This thesis represents one of the most thorough investigations into the application of titanium dioxide (TiO2) thin films to silicon photovoltaics (PV). TiO2 thin films have a long history of usage in the PV industry, primarily as an antireflection (AR) coating on screen-printed

(SP) solar cells. This is due, firstly, to the excellent optical properties of TiO2 and, secondly, to its low deposition cost.

In order to reduce the cost of fabricating buried-contact (BC) solar cells, an alternative dielectric thin film to silicon dioxide (SiO2) has to be found. The SiO2 growth step is lengthy and has to be performed at high-temperatures, increasing the complexity and cost of solar cell fabrication. Additionally, other research has demonstrated that aluminium alloying can occur in a phosphorus ambient,358 enabling two high-temperature steps to be combined into one. While silicon nitride has been demonstrated to be economic on a commercial scale for textured crystalline silicon (c-Si) wafers, TiO2 thin films appeared to be more suitable for solar cells fabricated on lower cost, multicrystalline silicon (mc-Si) wafers. This provided the motivation to investigate individual processes within the solar cell fabrication sequence and identify areas where TiO2 thin films could be used to either enhance performance and reduce fabrication costs.

Thin Film Requirements

Depending on the fabrication sequence of the BC solar cell, a single dielectric film, such as

TiO2, could serve multiple purposes:

• an AR coating,

• a dopant source for emitter diffusions,

199 200 8. Conclusions

• a film compatible with oxide passivation,

• a diffusion barrier to phosphorus diffusion,

• a mask for electroless metal plating.

These desired functions place relatively strict requirements on the dielectric film, demanding:

• a high refractive index and low extinction coefficient for excellent optical performance. This is especially important for mc-Si, for which no effective texturing method has been demonstrated on an industrial scale yet.

• a high chemical resistance to wet chemicals used in PV manufacturing. This requires a dense film with no pinholes to prevent chemicals and metal plating solutions from reaching to the front surface of the silicon wafer. Additionally, the films have to be insulating to prevent the electroless metal plating solution from adhering to the dielectric film.

• excellent stability during high-temperature processing in a variety of gas ambients,

including oxygen (O2), nitrogen (N2) and the dopant source, phosphorus oxychloride

(POCl3).

• and, to be attractive in a commercial environment, such a film needs to be deposited at a cost of less than US$0.05 per 5” square wafer.

While the BC technology was the main driver behind this work, much of the literature review is relevant to any work involving TiO2, while several of the experiments may be applicable to other solar cell designs such as those with evaporated or SP contacts.

Titanium Dioxide

TiO2 is a complex material with two of the crystalline phases – the metastable low- temperature phase of anatase as well as the stable high-temperature phase of rutile – as well as the amorphous phase being commonly observed in thin films. The transformation ◦ from amorphous TiO2 to anatase occurs at 300 − 350 C, while anatase can be transformed ◦ into rutile by heating at temperatures greater than about 800 C. TiO2 thin films used in the majority of PV applications are as deposited as amorphous films, due to the requirement of firing SP pastes through the layer in order to contact the solar cell emitter. The refractive index of these films is low (n ≈ 2.0 − 2.1atλ = 600 nm) and the chemical resistance to the majority of acids and bases is poor. Anatase films exhibit an increased refractive index of (n ≈ 2.3−2.45 at λ = 600 nm) and greatly enhanced chemical resistance. Rutile films exhibit refractive indices of n>2.45 at λ = 600 nm and also a significantly increased extinction coefficient due to an optical bandgap at about 3.05 eV. Therefore, anatase thin films seemed to afford the best optical properties and chemical resistance. 8.1 Summary 201

TiO2 Deposition Systems

Two systems were designed and fabricated by the author for the deposition of anatase thin films. Films deposited using a spray-deposition system exhibited ideal optical properties (high refractive index n and low extinction coefficient k) for a single-layer antireflection (SLAR) coating on a glass encapsulated solar cell. However, due to large variations in the film thickness, this system was converted into a simple chemical vapour deposition (CVD) system operating at atmospheric pressure. CVD-deposited TiO2 films exhibited much more uniformly thick layers, although the refractive index of these films was nearly 15% lower at the same deposition temperature of 450◦C.

No Contamination During High-Temperature Processing

Experimental results demonstrated that TiO2 is compatible with high-temperature process- ing without contaminating the wafers or furnaces. The bulk minority carrier lifetimes (τbulk) ◦ of TiO2-coated samples placed in a furnace for 2 hr at 950 C were maintained at greater than 2 ms. It was found that the spray-deposited TiO2 films are sensitive to the initial gas ambient in the furnace. Samples loaded in oxygen were stable, however TiO2 films loaded directly into a nitrogen ambient were reduced to a sub-oxide, most likely Ti2O3. The for- mation sub-oxide was predicted using thermochemistry analysis and confirmed using FTIR spectroscopy.

Overcoming the Surface Passivation Limitation

One major disadvantage in using TiO2 thin films in silicon solar cells is that they afford very little surface passivation to the silicon wafer. A novel method was developed by the author to overcome this limitation of achieving surface passivation on TiO2-coated silicon wafers. This involved growing a thin SiO2 layer at the TiO2:Si interface by oxidizing the wafer after TiO2

film deposition. The presence of the 6 nm-thick SiO2 layer was confirmed using scanning electron microscopy (SEM) images and X-ray photoelectron spectroscopy (XPS) analysis.

The increase in surface passivation afforded by the interfacial SiO2 layer results in a decrease in the emitter dark saturation current density (J0e) by nearly two orders of magnitude from −12 2 −14 2 ∼ 2 × 10 A/cm after TiO2 deposition to 4.7 − 7.7 × 10 A/cm . This demonstrates the ability of the TiO2/SiO2 AR coating to provide excellent surface passivation. The low J0e and high τbulk values demonstrated here are compatible with high-efficiency solar cells with an open circuit voltage (Voc) of the order of 700 mV.

TiO2 as a Phosphorus Diffusion Barrier

The results from this work indicate that a 70 nm thick TiO2 film can function as a phos- phorus diffusion barrier, thereby protecting the lightly-doped emitter (∼ 100 Ω/2)fromthe heavy phosphorus groove diffusion (< 5Ω/2). Although the TiO2 diffusion barrier resulted in only very light phosphorus diffusions underneath the film, a reaction in the POCl3 fur- nace substantially alters the optical and electrical properties and limits its usefulness. The 202 8. Conclusions

reaction between the TiO2 and phosphorus pentoxide (P2O5) led to the formation of new compounds on the wafer, determined to be either TiPxSiyOz or TiPxOy from Rutherford Backscattering (RBS) and thermochemistry analysis. Therefore, TiO2 must be described as a sacrificial diffusion barrier to phosphorus. This limitation means that in order for TiO2 to directly replace SiO2 in the BC process, a protective layer (such as SiO2)mustfirstbe applied on top of the TiO2 film.

TiO2 as a Phosphorus Dopant Source

By doping the TiO2 film with a small concentration of phosphorus, it was anticipated that these dopant atoms would diffuse out of the film and into the silicon during subsequent high- temperature processing. In this manner an n-type emitter could be created in the p-type silicon wafer, obviating the need for a separate emitter diffusion step. Results from this work have determined that the ability of TiO2 to act as a phosphorus dopant source is also somewhat limited. The increased conductivity of the TiO2 film due to the phosphorus incor- poration is not compatible with the insulating requirements of the BC solar cell electroless metal plating step. Also, it is anticipated that the increased optical absorption in the films, along with the lengthy diffusion times and temperatures involved (> 1hrat > 1000◦C) will limit the industrial uptake of this technology.

TiO2 Double-Layer Antireflection (DLAR) Coatings

The optical properties of TiO2 films have been varied greatly by altering the deposition and annealing conditions (temperature and ambient). The optical properties of the films were determined with high accuracy using spectroscopic ellipsometry over the wavelength range 350 − 1150 nm, at angles of 65 − 80◦. Optical modelling was performed using two main software packages, WVASE (J.A. Woollam) and TFCalc (Software Spectra). TiO2 films with refractive indices ranging from 1.726 to 2.633 at a wavelength of 600 nm have been produced. The low refractive index films were deposited in the presence of large amounts of water vapour, while the high refractive index films were achieved by annealing the as- deposited samples for a period of up to 6 hr at 1000◦C.

This variation in optical properties suggested that a DLAR coating design could be fabri- cated, with both layers being comprised of TiO2. A DLAR coating was designed for operation in air, with ideal refractive indices of nAR1 =1.58 and nAR2 =2.49, for the upper and lower layers respectively. The denser bottom TiO2 layer was obtained by annealing a TiO2 film ◦ deposited at 450 C, while a highly porous upper TiO2 layer was achieved by performing the deposition in the presence of water vapour. A distinctive double minima, characteristic of a DLAR coating, was observed in the reflectance spectrum. The best weighted average reflectance achieved was 6.54% for a planar crystalline silicon wafers and 8.04% for a planar multicrystalline silicon wafer (both in air). Both the individual layer thicknesses and the refractive indices were greater than ideal, indicating that further improvement is possible.

The real opportunity for a TiO2 DLAR coating comes, not for silicon solar cells in air, but for 8.2 Applicability to Various Solar Cell Processes 203 those encapsulated under ethyl vinyl acetate (EVA) and glass. The ideal refractive indices of nAR1 =2.07 and nAR2 =2.86 (at λ = 600 nm) are not achievable with any transparent materials, however modelling by the author has shown that a TiO2 DLAR coating with nAR1 =1.95 and nAR2 =2.60 could achieve a minimum weighted average reflectance (Rw) of 6.1%. With the inclusion of a 10 nm-thick SiO2 passivation layer at the interface, the Rw increased slightly to 6.9%. The performance of the DLAR coating was significantly higher than a SLAR, which exhibited an Rw of 8.3% and 8.1%, with and without the SiO2 layer, respectively. This indicates the opportunities for implementing such a DLAR coating as a cost-effective means for reducing the reflectance of mc-Si wafers.

Optical modelling determined that the majority of the absorptance in the short wavelength region occurred in the EVA layer, rather than in the glass or TiO2 layers. A PC1D simulation, which included absorption in all layers, of an encapsulated, planar BC solar cell with a TiO2

DLAR coating fabricated on low lifetime (τbulk = 100 µs) silicon resulted in a maximum short- 2 circuit current density (Jsc)of37.5 mA/cm . When a 10 nm-thick SiO2 passivation layer was 2 included at the TiO2:Si interface, this value decreased slightly to 37.2 mA/cm . The modelled efficiency for the latter solar cell was 18.6% (Voc = 640.6 mV), which is excellent considering the planar nature of the device. The fill-factor for this solar cell is FF =78.1%, which is achievable for a BC solar cell fabricated on mc-Si substrates. Use of the DLAR coating results in a 7% improvement in Jsc over a typical commercial TiO2 SLAR coating on the same device. For mc-Si wafers that have not received any form of gettering or hydrogenation, a bulk minority carrier lifetime of 20 µs is more typical. On this material, the simulated BC 2 solar cell with TiO2 DLAR coating achieved an efficiency of 17.4% (Jsc =35.8 mA/cm and

Voc = 625.0 mV).

8.2 Applicability to Various Solar Cell Processes

The application of TiO2 films to BC solar cells in order to reduce the fabrication costs and make was the BC technology more applicable to mc-Si wafers was the primary motivation of this research. However, the results of the experiments determining the plausible roles of

TiO2 films are relevant to other types of silicon solar cells. This section will highlight the opportunities for TiO2 thin films in BC, screen-printed (SP) and evaporated-contact solar cells.

BC Solar Cells

TiO2 is potentially a low-cost replacement for the thermally grown SiO2 layer as implemented in the original BC solar cell sequence. TiO2 can only be used as a direct replacement for

SiO2 layer in the standard BC processing sequence (refer to Section 1.4.3) if it is protected during the heavy phosphorus groove diffusion. Thus, the TiO2 film would function as a film compatible with surface passivation, an AR coating and an electroless metal plating mask. 204 8. Conclusions

The author demonstrated that a 200 nm-thick spin-on SiO2 layer is sufficient to protect the

TiO2 film during a phosphorus diffusion. This layer could remain on the solar cell until after electroless metal plating and subsequently be removed in hydrofluoric acid, which would not attack the TiO2. The cost using an additional protective barrier would be small, estimated at US$0.01 per 5” wafer, and along with the cost of depositing the TiO2 film (∼US$0.04 per 5” wafer) it is believed that this option is still cheaper than growing either SiO2 in an oxidation furnace or silicon nitride in an low-pressure chemical vapour deposition reactor.

A second option in the BC solar cell process is to use a single diffusion for both the emitter and grooves (typically 40−50 Ω/2). In this manner, the TiO2 is deposited prior to electroless metal plating, thereby avoiding exposing the film to phosphorus containing gas ambients.

The use of phosphorus doped titanium dioxide (TiO2:P) as a phosphorus dopant source is not applicable to the BC process, as the inclusion of P-dopant atoms makes the TiO2 film conductive and renders it useless as a electroless metal plating mask.

SP Solar Cells

The traditional role of TiO2 films in the PV industry is as an AR coating for SP solar cells.

An additional requirement of the TiO2 film in SP solar cells is that it “soft” enough for the SP paste to be able to be fired-through the layer and to contact the emitter. Deposition temperatures are typically about 200◦C, and the films retain many organics and exhibit a very low refractive index. However, during the brief firing step (∼ 800◦C for 1 min 30 s) the organics are baked off and the refractive index increases to about 2.25.

Thus, any role of TiO2 will only be successful if it is able to remain as an amorphous film up until the front contact firing step. Surface passivation is not usually a requirement for SP solar cells due to the heavily-doped emitter, which exhibits a phosphorus “dead-layer” and is insensitive to the front surface preparation conditions. A possible role of TiO2 in the

SP process would be the use of TiO2:P as an emitter dopant source. However, it may be the case that the times required for achieving phosphorus diffusion are not compatible with the metallisation paste firing times.

A TiO2 DLAR coating could be implemented if both of the layers were deposited as amor- phous and the front contact firing step was optimised to achieve high refractive indices. This may require the use of different pastes.

Evaporated-Contact Solar Cells

All of the novel roles of TiO2 may be applicable to a solar cell that uses evaporated contacts. There is greater flexibility in the design of such a solar cell due to the large number of steps (including photolithographic) involved. The most likely beneficial step would be the use of a

TiO2 DLAR coating, combined with the growth of SiO2 at the front surface for passivation purposes. The benefits of a DLAR coating for such a solar cell will be small, as these cells are normally fabricated on c-Si wafers, which are able to be chemically textured. However, the 8.3 Suggestions for Further Work 205

cost of a TiO2 DLAR coating will be substantially less than coatings comprised of evaporated films such as zinc sulphide and magnesium fluoride.

8.3 Suggestions for Further Work

There are several areas of research involving TiO2 that would appear to be promising. It should be noted that some of these “ideas” are merely the opinion of the author and there may be no scientific literature available to support these claims.

TiO2 would seem to be more compatible with boron than phosphorus. This has several implications, mainly for solar cells fabricated on n-type wafers. Firstly, shorter times and lower temperatures may be required in order to achieve a p-type emitter from boron-doped

TiO2. Secondly, it is possible that boron diffusion glass (B2O3) will not react with or increase the conductivity of the TiO2 film. If demonstrated to be true, this would make an n-type BC solar cell process very attractive. From the literature, it would also seem that tantalum pentoxide (Ta2O5) would be more compatible than TiO2 as a phosphorus dopant source.

Another interesting possibility is if (Al2O3) could act as a phosphorus diffusion barrier. If so, then an Al2O3/TiO2 DLAR coating could be used in the BC process.

The use of highly porous TiO2 films demonstrated in this work may be useful in other types of solar cells, such as the Gr¨atzel cell. In this cell, a thick (∼ 5 µm) TiO2 film is used a matrix, and flooded with electrolyte dye. A film with a high surface area is required so that the dye can contact as many points as possible.

A study of the optical and electrical properties of both doped TiO2 (say, with niobium) and reduced TiO2−x would be interesting. It may be possible to fabricate a transparent conduct- ing oxide layer from either of these films. TiO2-based films have the advantage that they can be deposited at high-speed, at atmospheric pressure and at low-cost. Therefore, it is possible that conducting layers for thin film modules could be deposited using atmospheric pressure chemical vapour deposition (APCVD), rather than more expensive sputtering techniques. 206 8. Conclusions Appendix A

TiO2 AR Coating Modelling Parameters

The fitting parameters used to model the dielectric constant function of the TiO2 films from

Sections 7.3 and 7.4 are included in this appendix. The fitting parameters, ε1(∞), Ai, Bi and

Eni are described in Section 4.6.3, along with Equation 4.8, which defines the mean-squared error (MSE).

In addition, the tabulated results of thickness and optical properties and their dispersive n and k spectra of some of the TiO2 thin films from Section 7.3 are also included in this appendix.

207 208 A. TiO2 AR Coating Modelling Parameters

A.1 Variation of n and k with Deposition Temperature

Table A.1: Fitting parameters used to model the dielectric constant (Lorentz double-oscillator model) of TiO2 films deposited using CVD with substrate temperatures 150 − 450◦C. The films are listed from lowest to highest deposition temperature with 50◦C intervals.

Sample ε1(∞) A1 B1 En1 A2 B2 En2 MSE Name (eV) (eV) (eV) (eV) TO-2-1 2.9145 0.27128 0.19185 3.5666 406.16 0.0044184 4.0648 1.977 TO-2-3 3.1056 0.19189 0.12474 3.5337 523.59 0.0023403 3.9785 1.480 TO-2-2 2.9594 0.33326 0.34003 3.5342 471.25 0.0036487 3.9302 2.900 TO-2-4 3.1135 0.27315 0.18417 3.5506 329.25 0.0061043 4.0383 1.801 TO-2-5 3.2556 0.26939 0.17283 3.5563 429.78 0.0054490 4.0787 1.328 TO-2-6 3.5291 0.30021 0.12739 3.5513 464.72 0.0047787 4.0308 1.154 TO-2-8 3.7805 0.56632 0.18329 3.5637 196.30 0.0083855 3.9464 1.084

Table A.2: Thickness, refractive index (at λ = 600 nm) and extinction co- efficient (at λ = 400 nm) of surface roughness and TiO2 layers deposited by CVD with substrate temperatures 150 − 450◦C.

Surface Layer TiO2 Layer T d nkd nk Sample dep surf TiO2 dsurf Name (◦C) (nm) (nm) dTiO2 TO-2-1 150 23.1 1.419 0.00149 87.0 1.879 0.00326 0.266 TO-2-3 200 26.7 1.419 0.00056 76.7 1.881 0.00121 0.348 TO-2-2 250 21.7 1.428 0.00721 102.0 1.901 0.01144 0.213 TO-2-4 300 14.8 1.451 0.00157 83.6 1.952 0.00343 0.177 TO-2-5 350 14.9 1.479 0.00133 76.2 2.012 0.00291 0.196 TO-2-6 400 29.6 1.506 0.00091 65.4 2.072 0.00199 0.453 TO-2-8 450 36.2 1.517 0.00204 63.0 2.097 0.00583 0.575 A.1 Variation of n and k with Deposition Temperature 209

Figure A.1: Dispersive relations for the (a) refractive index and (b) ex- tinction coefficient of the TiO2 layer of samples deposited using CVD at a range of deposition temperatures. 210 A. TiO2 AR Coating Modelling Parameters

Figure A.2: Dispersive relations for the (a) refractive index and (b) ex- tinction coefficient of the surface roughness layer of samples deposited using CVD at a range of deposition temperatures. A.1 Variation of n and k with Deposition Temperature 211

Table A.3: Fitting parameters used to model the dielectric constant (Lorentz double-oscillator model) of additional TiO2 films deposited us- ing CVD with substrate temperatures 250◦C (TO-3-8a) and 450◦C (TO- 3-1a).

Sample ε1(∞) A1 B1 En1 A2 B2 En2 MSE Name (eV) (eV) (eV) (eV) TO-3-8a 1.6877 73.911 0.03728 4.0662 175.45 0.33451 47.446 1.841 TO-3-1a 2.3540 7.5575 0.10686 3.6863 203.98 0.07535 7.3323 1.468

Table A.4: Thickness, refractive index (at λ = 600 nm) and extinction co- efficient (at λ = 400 nm) of surface roughness and TiO2 layers deposited by CVD with substrate temperatures of 250◦C and 450◦C.

Surface Layer TiO2 Layer T d nkd nk Sample dep surf TiO2 dsurf Name (◦C) (nm) (nm) dTiO2 TO-3-8a 250 14.1 1.455 0.00295 96.9 1.960 0.00975 0.146 TO-2-2 250 21.7 1.428 0.00721 102.0 1.901 0.01144 0.213 TO-3-1a 450 34.8 1.576 0.00719 56.0 2.225 0.01586 0.621 TO-2-8 450 36.2 1.517 0.00204 63.0 2.097 0.00583 0.575 212 A. TiO2 AR Coating Modelling Parameters

A.2 Variation of n and k with Annealing Temperature

Table A.5: Fitting parameters used to model the dielectric constant (Lorentz double-oscillator model) of TiO2 films deposited using CVD at ◦ ◦ 450 C and annealed at 450 − 1050 CinN2 for a period of 1 hr or 6 hr. MSE n 2 E 2 B (eV) (eV) 2 A n 1 E 1 B (eV) (eV) 1 A ( ∞ ) 1 ε 2 2 C) ann ◦ T hr Anneal in N hr Anneal in N Name ( Sample TO-4-1cTO-4-2c 500TO-4-3c 600 -1.8225 159.18TO-4-4c 700 -2.0668 183.72 1.6749TO-4-5c 800 -5.8686 330.64 1.5122 50.092TO-4-6c 900 -9.9736 0.71979 468.91 50.000TO-4-7c 28.747 950 -10.000 25.650 0.48890 1000 521.69TO-4-8c 41.470 -4.1817 0.074611 16.352 24.400 -5.2537 0.45637 1050 710.61 0.066487 3.9780 826.69 16.308 9.2176 -1.1817 0.15483 0.099716 2.033 4.0728 0.11261 118.49 12.512 9.2967 3.9666 0.13795 1.660 9.7629 0.39351 11.014 2.042 0.26935 7.9520 20.518 3.7119 0.28993 0.847 25.480 3.6038 0.23544 1.049 3.6601 0.17741 3.8497 1.819 1.477 3.8535 1.728 1 TO-3-1b 450 3.8177 7.5739 0.10541 3.5800 0.57569 0.73656 3.4834 3.434 6 TO-3-2bTO-4-1b 450TO-4-2b 500 -1.6873 151.30TO-4-3b 600 -1.8589 173.79 1.7879TO-4-4b 700 -2.6751 212.71 1.6131 50.160TO-4-5b 800 -7.3910 395.75 1.2638 50.014 10.840TO-4-6b 900 -10.000 0.59047 469.93 42.444 25.776TO-4-7b 950 -10.000 0.15430 21.082 0.48433 1000 498.96 34.738TO-4-8b 1.0791 0.055087 16.367 20.244 3.8313 -9.2070 0.44677 1050 1937.6 0.066355 3.8267 1.978 362.42 15.076 14.808 0.0091141 -10.000 0.06670 2.187 4.0231 0.41197 3.8748 383.08 13.109 36.479 0.24770 1.312 10.527 1.167 0.40191 16.847 0.21158 10.219 33.735 4.0054 1.108 19.637 3.6372 0.28769 0.096759 1.484 3.7903 3.7526 0.11758 1.384 1.824 3.7332 1.805 A.2 Variation of n and k with Annealing Temperature 213

Table A.6: Thickness, refractive index (at λ =600nm) and extinction coefficient (at λ =400nm) of surface roughness and TiO2 layers de- ◦ ◦ posited by CVD at Tdep = 450 C and annealed at 450 − 1050 CinN2 for 1 hr or 6 hr.

Surface Layer TiO2 Layer T d nkd nk Sample ann surf TiO2 dsurf Name (◦C) (nm) (nm) dTiO2

1 hr Anneal in N2 TO-3-1b 450 44.9 1.511 0.03861 80.4 2.082 0.08495 0.558 TO-4-1c 500 34.1 1.501 0.00639 65.4 2.061 0.01404 0.521 TO-4-2c 600 34.5 1.519 0.00582 61.0 2.100 0.01282 0.566 TO-4-3c 700 32.9 1.509 0.01122 68.4 2.077 0.02470 0.481 TO-4-4c 800 35.4 1.558 0.01905 64.2 2.184 0.04202 0.551 TO-4-5c 900 36.8 1.666 0.05753 63.2 2.423 0.12740 0.582 TO-4-6c 950 29.7 1.690 0.05737 60.6 2.477 0.12706 0.490 TO-4-7c 1000 14.2 1.722 0.04047 55.0 2.547 0.08962 0.258 TO-4-8c 1050 6.6 1.745 0.03775 63.5 2.598 0.08362 0.104

6 hr Anneal in N2 TO-3-2b 450 41.3 1.511 0.01326 70.5 2.081 0.02918 0.586 TO-4-1b 500 39.4 1.505 0.00495 67.5 2.069 0.01087 0.584 TO-4-2b 600 35.7 1.524 0.00583 60.0 2.110 0.01284 0.595 TO-4-3b 700 33.2 1.525 0.01155 66.0 2.113 0.02544 0.503 TO-4-4b 800 33.0 1.619 0.03084 59.4 2.319 0.06822 0.556 TO-4-5b 900 30.1 1.696 0.04645 63.7 2.489 0.10287 0.473 TO-4-6b 950 22.6 1.722 0.06018 63.7 2.546 0.13329 0.355 TO-4-7b 1000 4.7 1.747 0.02717 61.2 2.602 0.06017 0.077 TO-4-8b 1050 5.2 1.732 0.02966 66.5 2.569 0.06569 0.078 214 A. TiO2 AR Coating Modelling Parameters

Figure A.3: Dispersive relations for the (a) refractive index and (b) ex- ◦ tinction coefficient of the TiO2 layer of samples deposited at 450 C and annealed for 1 hr in N2. A.2 Variation of n and k with Annealing Temperature 215

Figure A.4: Dispersive relations for the (a) refractive index and (b) ex- ◦ tinction coefficient of the TiO2 layer of samples deposited at 450 C and annealed for 6 hr in N2. 216 A. TiO2 AR Coating Modelling Parameters

Figure A.5: Dispersive relations for the (a) refractive index and (b) ex- tinction coefficient of the surface roughness layer of samples deposited at ◦ 450 C and annealed for 1 hr in N2. A.2 Variation of n and k with Annealing Temperature 217

Figure A.6: Dispersive relations for the (a) refractive index and (b) ex- tinction coefficient of the surface roughness layer of samples deposited at ◦ 450 C and annealed for 6 hr in N2. 218 A. TiO2 AR Coating Modelling Parameters

A.3 Variation of n and k with Deposition Ambient

Figure A.7: Dispersive relations for the (a) refractive index and (b) ex- ◦ tinction coefficient of the TiO2 layer of samples deposited at 250 C and ◦ 450 C in the presence of H2O vapour. Some samples were annealed at ◦ ◦ either 450 Cor1000 CinN2 for a period of 1 hr or 6 hr. A.3 Variation of n and k with Deposition Ambient 219

Table A.7: Fitting parameters used to model the dielectric constant (Lorentz double-oscillator model) of TiO2 films deposited using CVD at ◦ ◦ 250 C and 450 C with H2O vapour. Some samples were annealed at ei- ◦ ◦ ther 450 Cor1000 CinN2 for a period of 1 hr or 6 hr. MSE n 2 E 2 B in the same high- (eV) (eV) 2 :Si interface. 2 2 A at the TiO 2 n 1 E 1 . 1nmofSiO B and subequently 5 hr 55 min in N 2 1 A C) (hr) (eV) (eV) ( ∞ ) ◦ 1 ε 2.1066 13.863 4.9956 86.830 232.99 0.006443 4.3773 2.456 0.86073 1.1237 1.9594 9.1427 594.07 0.029195 6.6227 1.754 0.65607 6.0971 0.39140 3.6764 172.98 0.95360 57.808 6.552 † C) ( ann − − ◦ T 2 2 C) ( dep ◦ T Sample TO-3-5b was annealed for 5 min in O † temperature step. This resulted in the growth of 11 hr Anneal in N hr Anneal in N Name ( Sample TO-3-7c 250 1000 0.10645 10.537 6.3942 25.798 3.6709 0.18753 3.7506 3.112 TO-3-4a 450 TO-3-6a 250 TO-3-5a 450 1000 -0.97474 6.7456 0.29274 3.7236 323.72 0.37816 24.545 2.002 1 TO-3-5b 450 1000 No Anneal TO-3-6b 2506 450 1.9741 19.200 3.7608 100.00 246.31 0.003356 3.9991 4.900 220 A. TiO2 AR Coating Modelling Parameters

A.4 Variation of n and k with Annealing Ambient

Table A.8: Fitting parameters used to model the dielectric constant (Lorentz double-oscillator model) of TiO2 films deposited using CVD at 450◦C. Some samples received a 5 min oxidation prior to all samples being ◦ annealed at 1000 CinN2 for a total period of 1 hr or 6 hr. MSE n 2 E 2 B (eV) (eV) 2 A O vapour. 2 n 1 E 1 B (eV) (eV) 1 A ( ∞ ) 1 ε -5.2537-5.4320 826.69 0.11261 64.836 9.7629 0.12730 20.518 3.9827 0.23544 34.292 3.8497 5.0079 1.477 19.722 1.950 -3.6533-1.1366 30.090 0.12581 40.871 3.8632 0.18648 301.74 3.8823 0.21910 241.44 7.7274 0.30138 1.083 13.965 1.896 0.65607 6.0971 0.39140 3.6764 172.98 0.95360 57.808 6.552 -0.97474 6.7456 0.29274 3.7236 323.72 0.37816 24.545 2.002 ? 2 ✗ ✗ ✗ ✓ ✓ ✓ 5 min † † C) O dep ◦ T Samples TO-3-5a and TO-3-5b were deposited in the presence of H † hr Anneal hr Anneal Name ( Sample TO-4-7c 450 TO-3-3c 450 TO-3-3aTO-3-5a 450 450 1 TO-3-5b 450 6 TO-3-3b 450 A.4 Variation of n and k with Annealing Ambient 221

Figure A.8: Dispersive relations for the (a) refractive index and (b) ex- ◦ tinction coefficient of the TiO2 layer of samples deposited at 450 C. Some samples received a 5 min oxidation prior to all samples being annealed at ◦ 1000 CinN2 for a total period of 1 hr or 6 hr. 222 A. TiO2 AR Coating Modelling Parameters

Table A.9: Fitting parameters used to model the dielectric constant (Lorentz double-oscillator model) of TiO2 films deposited using CVD at 450◦C and oxidised for a period of 1 − 16 min at a loading temperature of 800◦C with the furnace set to ramp to 1000◦C. MSE n 2 E 2 B 2 A n 1 E 1 B 1 A ( ∞ ) 1 ε ann t Name (min) (eV) (eV) (eV) (eV) Sample TO-4-1aTO-4-2a 1TO-4-3a 2TO-4-4a -2.0959 4TO-4-5a -1.8936 201.43 8 16 -1.9082 215.81 1.3653 -3.4864 190.35 1.3091 50.705 0.737663 23.324 208.43 1.3604 13.996 52.835 0.10798 0.27947 24.777 1.1598 49.804 3.9790 3.6323 0.11224 28.435 35.499 1.760 209.27 4.0005 0.12884 24.532 0.87688 1.779 4.0914 0.22635 50.000 1.524 4.0323 2.393 1.536 A.4 Variation of n and k with Annealing Ambient 223

Table A.10: Thickness, refractive index (at λ =600nm) and extinction coefficient (at λ =400nm) of surface roughness and TiO2 layers de- posited using CVD at 450◦C and oxidised for a periof of 1 − 16 min at a loading temperature of 800◦C with the furnace set to ramp to 1000◦C

Surface Layer TiO2 Layer t d d nkd nk Sample ann SiO2 surf TiO2 dsurf Name (min) (nm) (nm) (nm) dTiO2 TO-4-1a 1 0 32.4 1.497 0.00974 72.9 2.050 0.02142 0.444 TO-4-2a 2 0 29.3 1.519 0.01017 61.7 2.100 0.02240 0.475 TO-4-3a 4 0.4 32.0 1.529 0.01232 65.6 2.122 0.02716 0.488 TO-4-4a 8 3.7 28.2 1.602 0.03239 57.7 2.283 0.07163 0.489 TO-4-5a 16 8.7 33.7 1.678 0.07174 60.8 2.449 0.15890 0.554 224 A. TiO2 AR Coating Modelling Parameters

Figure A.9: Dispersive relations for the (a) refractive index and (b) ex- ◦ tinction coefficient of the TiO2 layer of samples deposited at 450 C and oxidised for a periof of 1 − 16 min at a loading temperature of 800◦C with the furnace set to ramp to 1000◦C. A.5 TiO2 DLAR Coatings 225

A.5 TiO2 DLAR Coatings

Table A.11: Fitting parameters used to model the dielectric constant (Lorentz double-oscillator model) of TiO2 DLAR coatings measured using spectroscopic ellipsometry. A dash (−) indicates that the second oscilla- tor was not required in the model and that the amplitude was zero. MSE 3.643 3.945 3.384 n 2 E 2 B (eV) (eV) 2 −−− −−− −−− A n 1 E 1 B (eV) (eV) 1 A ( ∞ ) 1 ε High 3.6191 3969.7 0.0022149 4.2949 3.7625 0.26300 3.4587 High 4.6997 2907.7 0.0015748 3.8822 4.8528 0.28336 3.4819 High 5.6057 2370.6 0.0015067 3.7147 4.6486 0.19178 3.4987 Name Index Sample High/Low TO-5-5d Low 1.0261 17.388 5.0788 82.390 TO-5-6d Low 1.2686 19.988 5.0002 74.719 TO-5-8d Low 0.30975 37.791 5.0586 84.695 226 A. TiO2 AR Coating Modelling Parameters Bibliography

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