Cathode Interface Structure in Organic Semiconductor ... · Cathode Interface Structure in Organic Semiconductor Devices by Ayse Zeren Turak A thesis submitted in conformity with

Cathode Interface Structure in Organic Semiconductor Devices

by

Ayse Zeren Turak

A thesis subm itted in conform ity w ith the requirem ents for

the degree of Doctor of Philosophy

G raduate D epartm ent of M aterials Science and Engineering

U niversity of Toronto

C opyright © by Ayse Turak 2006

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CanadaReproduced with permission of the copyright owner. Further reproduction prohibited without permission.

AbstractCathode interface structure in organic sem iconductor devices

Ayse Zeren Turak

Doctor o f Philosophy

Graduate Department of Materials Science and Engineering

University o f Toronto

2006

As organic semiconductor technology matures, enhancement requires

understanding/engineering o f the cathode/organic interface. In this work, using X-ray

photoelectron spectroscopy (XPS) and common materials for organic light emitting diodes

(OLEDs), the expected interfacial structure in conventionally fabricated devices has been

described and some simple predictive methods developed.

The buried electrode/active layer interface was examined by analysing: 1. both sides

o f the interface in conventionally fabricated devices under high vacuum with the unique peel-

off technique, and 2. monolayers of one material grown atop another. Connections were

drawn between the interfacial structures in devices, those observed during traditional surface

science investigations, and the device behaviour. A critical insight is that no one metal or

metal/interlayer combination may be used as a universal cathode. Rather, certain criteria for

interfacial structure and stability must be confirmed to ensure adequate performance. This

can be determined through simple material property information, such as lattice constants, or

with inorganic analogues for organic molecules.

For combinations of metals and 8 -tris(hydroxyquinoline aluminum) (Alq3),

interfacial reactions can be predicted by assuming AI2O3 as an inorganic analogue. Using this

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analogue, molecular fragmentation may be described as a simple metal-exchange oxidation-

reduction reaction.

As cathode complexity increases, such simple descriptions lose validity. This work

shows that all three components (organic/LiF/metal) are required to adequately describe the

interfacial structure of bi-layer cathodes. The major conclusions regarding the role of LiF are:

• that 5-10A LiF changes the cathode oxidation behaviour, predicted by the lattice

mismatch of the interlayer with the metal. Oxidation is suppressed for Al, which is well

matched to LiF; for Mg, which has poor matching, preferential formation of carbonates

occurs. Device behaviour is related to the metal oxidation, such that Al/LiF cathodes are

superior to Mg/LiF ones.

• that near the interface, LiF forms charge transfer complexes with electron transporting

molecules.

• that the cathode should be considered a metal-insulator-metal capacitor with the organic

layer acting as the bottom electrode. The usable thickness of LiF is dependent on the

conductivity of the layer.

These insights indicate some of the conditions necessary for adequate device

performance and longevity, useful for future device optimization.

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Acknowledgements

First and foremost, I would like to express my sincere thanks to my supervisor

Professor Zheng-Hong Lu, for his support and guidance throughout this project.

I would also like to thank the members of my reading committee, who all provided

much insight and support and lively discussion over the years of this project. My thanks go to

Professor Iain Sommerville, Professor Charles Mims, and Professor Bob Pilliar for all of

their constructive advice and encouragement.

I would also like to thank all the members o f the Lu group, past and present, who

have been excellent colleagues and good friends over the last five years. Especially, my

gratitude to Dr. Daniel Grozea, with whom I have worked very closely over many years,

right from the beginnings when the current lab was nothing but an old machine shop. For

excellent discussions, for some sample preparation, and for general camaraderie, I would also

like to thank Drs. Changjun Huang and Sijin Han.

I would like to thank our colleagues at the National Research Council Institute for

Microstructural Studies - Drs. Chandra Dharma-Wardana and Marek Ziegrzgi for theoretical

density functional calculations, and Jeff Fraser for scanning electron microscopy.

My thanks also go to Dr. Bradley Diak at Queen’s University, for supplying some

substrates at a critical junction, and for general support and mentoring over the years.

I am very grateful to everyone at the Department of Materials Science and

Engineering over the years, especially Rob Guzzo and Phil Egberts for always being willing

to take a break, Sal Boccia for always having just the right piece of equipment; and Louisa,

Teresa, and Fanny for always having time for a quick question.

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None of this would have been possible without the constant love and support o f my

family. My parents, A1 and Nel, and my brother, Devrim, have always been the support

beneath my every success - encouraging me, cheering me up, keeping me motivated, even

reading my final drafts. For always keeping me grounded in my greatest ambitions, I dedicate

this thesis to them.

Finally, I would like to acknowledge the Natural Sciences and Engineering Research

Council of Canada, Materials and Manufacturing Ontario (currently Ontario Centres of

Excellence), and the Ontario Graduate Scholarship for their generous financial support.

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Contents

List of Figures ..................................................................................................... x

List of Tables ............................................................................................... xviii

Nomenclature ................................................................................................... xx

Chapter 1 Introduction.........................................................................................11.1 Interfaces in organic electronics........................................................................................11.2 Thesis organization.............................................................................................................41.3 References..........................................................................................................................5

Chapter 2 OLED fundamentals..........................................................................62.1 OLEDs and organic conductors.........................................................................................6

2.1.1 Device operation........................................................................................................... 62.1.2 Device structures........................................................................................................... 82.1.3 Organic electron conduction layers:...........................................................................9

2.2 Role of the interface in OLEDs..................................................................................... 142.2.1 Injection........................................................................................................................142.2.2 Device reliability......................................................................................................... 15

2.3 Cathode performance........................................................................................................162.3.1 Elemental metal cathodes.......................................................................................... 162.3.2 Alloy cathodes............................................................................................................ 172.3.3 Bi-layer cathodes........................................................................................................ 19

2.4 References....................................................................................................................... 22

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Chapter 3 Theoretical background of X-ray photoelectronspectroscopy and its applicability to interfacial analysis in OLEDs................................................................................................ 25

3.1 Basic principles.................................................................................................................25

3.2 Chemical shift................................................................................................................... 283.2.1 Prediction of chemical shift from absolute binding energy calculations 283.2.2 Extension of Seigbahn theory....................................................................................30

3.3 Use of secondary effects for analysis............................................................................323.3.1 Shake up features.........................................................................................................323.3.2 Auger excitation...........................................................................................................35

3.4 Charging in X-ray photoelectron spectroscopy........................................................... 383.4.1 Charge compensation.................................................................................................. 413.4.2 Use o f charging for electrical information with X P S ............................................ 43

3.5 Angle resolved XPS......................................................................................................... 443.5.1 Information depth of photoelectrons.........................................................................443.5.2 Thickness and coverage dependence of overlayers................................................46

3.6 Equilibrium chemical states analysis o f interfaces in OLEDs with X PS ................. 483.6.1 Low work function metal cathodes........................................................................... 493.6.2 Bilayer cathodes...........................................................................................................543.6.3 Limitations of previous studies utilizing X PS.........................................................56

3.7 References........................................................................................................................ 58

Chapter 4 Experimental..................................................................................... 624.1 Molecular beam deposition/Vapour phase deposition theory.................................... 62

4.2 Instruments....................................................................................................................... 674.2.1 MAC in-situ system.................................................................................................... 674.2.2 Cluster tool................................................................................................................... 74

4.3 In-situ peel off m ethod.................................................................................................... 764.4 Other analysis techniques................................................................................................794.5 References........................................................................................................................ 79

Chapter 5 Metal/Alq3 interface structures....................................................815.1 Introduction...................................................................................................................... 815.2 Experimental..................................................................................................................... 825.3 Results and discussion..................................................................................................... 835.4 Summary............................................................................................................................935.5 References........................................................................................................................ 93

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Chapter 6 LiF/metal bilayer structures I - Case of Al/LiF.........................956.1 Introduction.......................................................................................................................95

6.2 Experimental..................................................................................................................... 976.2.1 Sample preparation and analysis...............................................................................976.2.2 Thickness and coverage determination....................................................................98

6.3 Oxidation and surface structure of A1 surfaces..........................................................1016.3.1 Oxidation products....................................................................................................1016.3.2 Surface oxidation kinetics...................................................................................... 1036.3.3 Surface oxide structure............................................................................................ 1096.3.4 Impact o f LiF on metal surface oxidation..............................................................I l l

6.4 Estimation o f device failure due to oxidation of Al/LiF based cathodes................. 1126.5 Interfacial chemical structure at the Al/LiF/organic interface.................................. 1176 .6 Summary.......................................................................................................................... 1216.7 References.......................................................................................................................122

Chapter 7 LiF/metal bilayer structures II - Case of Mg/LiF................... 1247.1 Introduction.....................................................................................................................1247.2 Experimental....................................................................................................................1267.3 Oxidation products and kinetics of Mg surfaces........................................................ 1267.4 Chemical structure at organic interface with Mg/LiF in device structures 1337.5 Summary.......................................................................................................................... 1407.6 References.......................................................................................................................141

Chapter 8 LiF interaction with organics...................................................... 1438.1 Chemical structure of Al/LiF/Alq3 in organic light-emitting diodes........................143

8.1.1 Introduction................................................................................................................1438.1.2 Experimental............................................................................................................. 1458.1.3 Results and discussion............................................................................................. 1468.1.4 Summary.................................................................................................................... 151

8.2 LiF interaction with C6o................................................................................................. 1528.2.1 Introduction................................................................................................................1528.2.2 Experimental............................................................................................................. 153

8.2.3 Results and discussion............................................................................................. 1548.2.3.1 F Is core level for Cgo-LiF interaction........................................................1548 .2.3.2 Geometry optimized structures and theoretical prediction of the F Is

core level shift...............................................................................................1568 .2.3.3 C Is shake-up satellites for deposited monolayers.................................... 1618.2.3.4 Theoretical support for LiF**C6o complex formation................................ 1668.2.3.5 Growth morphology and critical thickness for F Is peak appearance... 1698.2.3.6 Other LiF/organic interactions.....................................................................172

8.2.4 Summary....................................................................................................................176

8.3 References.......................................................................................................................177

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Chapter 9 LiF layer properties........................................................................ 1819.1 Introduction.....................................................................................................................181

9.1.1 Device behaviour....................................................................................................... 182

9.2 Experimental....................................................................................................................183

9.3 Results and discussion................................................................................................... 1849.3.2 LiF growth on organic surfaces............................................................................... 184

9.3.3 Resistivity effects as observed by XP S ..................................................................1919.3.3.1 Charging effects in XP S ................................................................................ 191

9.3.4 X-ray induced charging effects for LiF coated films..............................................1929.3.4.1 Transient effects............................................................................................. 1959.3.4.2 Estimation of film conductivity from transient effects.............................. 197

9.4 Summary.......................................................................................................................... 207

9.5 References.......................................................................................................................208

Chapter 10 Interfacial structure models and conclusions...........................21010.1 Introduction.................................................................................................................... 210

10.2 Metal/Alq3 interfaces..................................................................................................... 212

10.3 LiF as an interlayer.........................................................................................................21410.3.1 LiF impact on the cathode metal............................................................................. 215

10.3.1.1 A l/LiF.............................................................................................................. 21510.3.1.2 Mg/LiF............................................................................................................. 21610.3.1.3 Lattice constants as a predictive tool............................................................217

10.3.2 LiF impact on the organic........................................................................................ 21910.3.3 LiF interlayer properties.......................................................................................... 220

10.4 Metal/LiF/organic system..............................................................................................221

10.5 Cathode selection for organic electronics.................................................................... 222

10.6 Future work......................................................................................................................224

10.7 References.......................................................................................................................228

Appendix A List of empirical charge-binding energy............................... 229

Appendix B Schematic of OMAC................................................................... 233

Appendix C Data analysis in XPS ................................................................. 235

Appendix D Equations for quantitative X PS...............................................241

Appendix E Structure calculations for C6o-LiF interaction.......................246

Appendix F Summary of observed F Is core level in all experiments... 249

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List of Figures

Figure 2-1 Schematic of typical OLED structure........................................................................... 9

Figure 2-2 Alq3 molecule (a) planar structure (b) Three-dimensional model o f themeridinal isomer [11]...................................................................................................................10

Figure 2-3 Ceo molecule....................................................................................................................11

Figure 2-4 Monte Carlo simulations o f C6o growth [20] on (a) Si and (b) diamond substrates (c) low density films grown by throwing C60 molecules at Si substrates at low energy..................................................................................................................................... 12

Figure 2-5 X-ray scan o f a 4450A thick C6o film deposited onto an off-axis cut singlecrystal sapphire substrate. The markers indicate the calculated diffraction lines from a face-centred-cubic cell seen in diffraction from bulk C6o powder [22]..............................12

Figure 2-6 Various substrate/molecule interactions that can lead to dipole formation [39] ..15

Figure 2-7 Device performance as a function of the cathode metal work function (a) relative luminance at a constant current density (b) relative efficiency at a constant luminance [54]............................................................................................................................. 17

Figure 2-8 Effect of interlayer with a variety of cathodes (a) for current density and (b) for relative efficiency for a LiF interlayer from Stofiel et al. [9], (c) for MgO and GeCL interlayers from Hung et al. [36].....................................................................................21

Figure 3-1 Photoemission process ........................................................................................26

Figure 3-2 Relation between the energy levels in a solid and the electron-energy distribution in the photo-emitted spectrum. The excitation source determines the range of interest. [2].................................................................................................................... 27

Figure 3-3 Core levels for various elements [2]..........................................................................27

Figure 3-4 Orbital redistribution due to the formation of a core ho le ....................................... 33

Figure 3-5 Schematic o f the XPS spectrum under the sudden approximation........................ 34

Figure 3-6 Auger emission process [after 2 1 ].............................................................................35

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Figure 3-7 Wagner plot for copper showing the Cu 2pm binding energy and Cu L 3M 4 5M 4 5

Auger kinetic energy for different chemical states. The straight lines with slope -1 represent compounds with the same Auger parameter, while those with slope -3 represent those with the same initial state effects. The binding energy and the Auger kinetic energy are referenced to the adventitious C Is line, set at 284.8eV [22].................. 37

Figure 3-8 (a) Schematic o f charging inside a semiconductor or insulator during XPS measurement (b) Binding energy modification due to positive charging in the sample electronically decoupled from the spectrometer (adapted from [27]).................................39

Figure 3-9 Sources o f charge compensation. Left: sources issued from the specimen holder. Right: sources issued from the surroundings o f the specimen that are normally to ground. The dotted lines indicate the incoming X-rays, and the solid lines the electron movement in the sample (adapted from [29])..................................................... 43

Figure 3-10 Schematic of angle resolution.................................................................................... 44

Figure 3-11 Exponential decay of lossless electron escape with depth o f creation o fphotoelectron [39]........................................................................................................................45

Figure 3-12 Enhancement of surface composition in core level intensity at grazingangles for Si 2p [21].................................................................................................................... 46

Figure 3-13 Response of various overlayer configurations to changes in the take-offangle for photoelectrons..............................................................................................................47

Figure 3-14 Deconvoluted O Is core level [53]........................................................................... 50

Figure 3-15 N Is evolution with K deposition [60].....................................................................51

Figure 3-19 UPS spectrum for the HOMO region of Alq3 after deposition o f variouscathodes [52]................................................................................................................................ 53

Figure 4-1 Multi-Access Chamber (MAC) System......................................................................68

Figure 4-2 Resistive sources used for thermal evaporation of inorganic materials in theOMAC chamber [1 ].................................................................................................................... 70

Figure 4-3 Schematic of cathode thermal evaporation source (a) side/front view(b) top view showing shielding and crucible configuration.................................................. 70

Figure 4-4 Kurt J. Lesker OLED cluster tool in the clean room ................................................ 75

Figure 4-5 Schematic of peel-off procedure. Top panel shows the removal o f the glass substrate to expose the cathode surface for analysis. Bottom panel shows the removal of the cathode layer to expose the organic surface for analysis indicated by the circled areas. Conductive carbon tape was used in both instances to adhere the sample to the sample holder to minimize any charging effects. The cleavage plane during peel-off is indicated by the heavy dotted line in the “side view” section............................................77

Figure 4-6 Sample holder schematic for substrate scoring before peel-off...............................78

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Figure 5-1 Top panel shows Al 2p core level spectra recorded on various as peeled off metal surfaces: Ag surface shown as circles, Mg surface shown as open triangles and Mg:Ag alloy surface shown as solid circles. The bottom panel shows curve fitting results o f Al 2p recorded on the Mg:Ag surface. The experimental data (solid circles) can be well fitted by the sum (solid line) of two separate spin-orbit doublet peaks (dashed lines), one metallic state at 72.7 eV and another Al3+ state at 74.4 eV...................85

Figure 5-2 XPS depth profile o f as-recorded Al 2p core levels obtained from: (a) Agcathode, (b) Mg cathode and (c) Mg:Ag alloy cathode...........................................................86

Figure 5-3 XPS depth profile o f intensity normalized Mg 2p core levels obtained from:(a) Mg cathode and (b) Mg:Ag alloy cathode...........................................................................88

Figure 5-4Schematic summary of various interface structures: (a) Mg: Ag/Alq3 interface,(b) Mg/Alq3 interface, (c) Ag/Alq3 interface, and (d) Au/Alq3 interface..............................88

Figure 6-1 Al 2p core levels for uncoated Al and 10A LiF coated Al for exposure times of (a) 25 mins and (b) 1500 hrs. Due to the insulating nature of LiF, the coated surface shows an increasing surface charging effect with time of 0.06 eV and 0.47eV for (a) and (b) respectively........................................................................................................102

Figure 6-2 Growth of oxide on Al surfaces, monitored by XPS, for thickness as estimated by the simple overlayer model. Lines represent a linear sum of reduced squares best fit o f the data for the uncoated and 10A LiF coated substrates.Uncoated and 5A LiF coated Al both show a bend in the curve at around 60 hrs.The open triangles represent the predicted oxide values scaled by the LiF coverage as predicted by ARXPS.............................................................................................................104

Figure 6-3 Mott-Cabrera oxidation behaviour for uncoated and 10A LiF coated Al surfaces. The solid lines represent a linear sum of reduced squares best fit of the data substrates.............................................................................................................................107

Figure 6-4 Determination o f LiF coverage using the simple patchy overlayer model (equation 6-2). The various lines represent different values of the coverage and LiF thickness, which were the only variables used to fit the data. The close up section on the right hand side shows the predicted angular dependence with a 5A LiF layer at different coverages. The Levenberg-Marquardt reduced chi squared fit of the experimental data (dotted line) indicates coverage o f 15%..................................................109

Figure 6-5 Structure model comparisons for LiF coated Al surfaces for exposure times of (a) 25 mins with 5 A LiF coverage and (b) 1500hrs exposure with 10A LiF coverage. Lines represent Levenberg-Marquardt reduced chi squared fit of the experimental data for various structure models. The solid line represents an embeddedstructure, the dashed line a columnar structure, and the dotted line a multilayer structure assuming a complete LiF layer at the interface......................................................110

Figure 6-6 Schematic oxide growth model on LiF coated Al surfaces, (a) Initially, growth occurs between LiF islands, producing a columnar structure. As growth progresses, Al diffuses through the LiF islands and growth occurs over the islands, leading to (b) an embedded structure...................................................................................... I l l

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Figure 6-7 Comparison of device behaviour on the first and 7th (final) day o f the experiment. The circles and triangles indicate the behaviour of stressed devices and unstressed devices after the same length o f exposure, indicating similar behaviour.The behaviour of the stressed device with 20A LiF no longer shows appropriatediode behaviour after the 2nd day of stressing, but the unstressed device on the finalday indicates a similar trend as for all the other thicknesses................................................ 113

Figure 6-8 (a) Current decay measurement for C6o based devices with Al/LiF cathodes o f varying LiF thickness. In region (1), the performance decays in relation to the LiF thickness. After one day of exposure, region (2), the device performance decays exponentially, with the same decay constant. The solid lines are guides to the eye, but in region (2) indicate the reduced squares best fit o f exponential decay with a decay constant as determined from figure 6-9. The device with a 20A LiF layer has very similar exponential decay behaviour to that of the 30A LiF device,, (b) Thethickness dependent percentage decrease in current after the first day............................... 114

Figure 6-9 Renormalized maximum current achievable over time. The solid linerepresents a linear sum of reduced squares best fit o f the data.............................................115

Figure 6-10 The maximum decay time to reach 10% of initial device performance.The lines are just a guide to the eye......................................................................................... 116

Figure 6-11 Al 2p core level for (a) Al surface (b) A1/100A LiF surface after peel-off at the cathode/organic interface (c) the metal surface o f the cathode after Ar+ sputtering (d) the sputter profile through the thickness of the LiF layer for A1/100A LiF cathodes showing the evolution of the chemisorbed Al.................................................118

Figure 6-12 I-V characteristics o f the C6o sandwich diodes with Al and Al/LiF after exposure to air for 1 hr and then baked in vacuum for 24 hrs (excerpted with permission from [20], Copyright 2005, American Institute of Physics)............................. 120

Figure 6-13 Depth profile results for a 200A LiF layer, showing the complete blocking o f oxygen diffusion from the organic layer. The residual Alq3 is removed after the first two cycles. Alq3 shows very little lateral diffusion of oxygen, so oxidation of metal surface due to diffusion from outer cathode surface through grain boundaries or during initial deposition. [42]...................................................................................... 120

Figure 7-1 Mg 2p core level for uncoated Mg and 10A LiF coated Mg for exposures times of (a) 7.8hrs and (b) 1500hrs. Both uncoated and coated surfaces show a pronounced high binding energy shoulder, corresponding to a superposition of Mg(OH)2 and MgO states. For the LiF coated surface, there is a shift o f 0.1 and 0.47eV due to surface charging for (a) and (b) respectively.................................................127

Figure 7-2 (a) The binding energy difference between the most intense peak from metallic Mg and that from the higher binding energy side of the Mg 2p core level.Open diamonds represent the shift in the hydroxide binding energy from charging due to the presence of LiF as deduced from the shift to the Al oxide peaks in chapter 6. Lines are just a guide to the eye (b) The change in the FWHM of the high binding energy component of the Mg 2p core level. The lines represent a linear sum of reduced squares best fit of the data.......................................................................................... 128

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Figure 7-3 (a) Curve fitting results for Mg 2p o f 10 A LiF coated Mg at 1500 hrs exposure. The experimental data (open diamonds) can be well fitted by the sum (solid line) of three separate peaks (dashed lines), one metallic state at 49.5eV, one hydroxide/oxide state at 51eV, and a carbonate state at 52eV. Oxide values include a 0.47eV charging offset due to the insulating nature o f LiF on the surface of Mg.(b) C Is core level of 10A LiF coated Mg (open diamonds) surfaces after 1500 hrs exposure, with three chemical states attributable to adventitious C (284.6eV and 286eV) and MgC0 3 (289.5eV). There may also be a slight contribution at 291eV, also likely due to adventitious C. (c) Curve fitting results for O Is o f 10 A LiF coatedMg at 1500 hrs exposure. The experimental data (open diamonds) can be well fitted by the sum (solid line) o f three separate peaks (dashed lines), an oxide state at 531.3eV, a hydroxide state at 532.9eV, and a carbonate state at 533.9eV. There is likely also a contribution from the adventitious C-OH beneath the carbonate peak at 531.3eV that could not be resolved......................................................................................129

Figure 7-4 Growth of oxide on Mg surface monitored by XPS (a) for uncoated Mg and 10A LiF coated Mg surfaces. Lines represent a linear sum of reduced squares best fit of the data. 10A LiF coated Mg shows abend in the curve after 100 hrs. (b) Growth of various oxide components for the LiF coated surface. The onset o f the bend observed in (a) corresponds to a shift from carbonate dominated growth to hydroxide dominated growth. The dotted lines are just a guide to the eye........................................... 131

Figure 7-5 O Is core level for 10 A coated and uncoated Mg surface after 1500 hrs exposure, indicating a greater amount of unconverted MgO for uncoated Mg. For the LiF coated surface, there is a shift o f 0.47 eV due to surface charging........................132

Figure 7-6 Al 2p core level recorded for the Mg/LiF surface. The experimental data can be well fitted by the sum (solid line) of two separate peaks (dashed lines), one metallic state at 72.9eV and another Al3+ state at 74.4eV.................................................... 134

Figure 7-7 Mg 2p core level for both Mg and Mg/LiF cathodes at the cathode side of the as-peeled interface, indicating (a) the difference in the high binding energy shoulder for the two surfaces of 0.5eV. (b) and (c) The curve fitting results of Mg 2p recorded on the Mg and Mg/LiF surface respectively. The experimental data (solid circles) can be well fitted by the sum (solid line) o f separate peaks (dashed lines). In both cases, the metallic state is at 48.5eV. For (b) the oxide peak corresponds to hydroxide formation at 1.4eV above the metallic. The two peaks in (c) are located at2.1 and 3.5eV above the metallic peak ..... 136

Figure 7-8 Potential reaction products formed at the cathode/Alq3 interface for Mgcathodes.......................................................................................................................................137

Figure 7-9 (a) Luminance-voltage characteristics for Mg cathode devices with and without a 10A LiF interlayer (b) Current-voltage characteristics adapted from M.StoBel etal. [2]...........................................................................................................................140

Figure 8-1 Various XPS core level spectra recorded on the organic side of the cleavedcathode/organic interface.......................................................................................................... 146

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Figure 8-2 F Is core level spectra recorded on the organic side of the cleaved interface with LiF interlayer thickness of: (a) 3 A, (b) 15A, and (c) 200A, respectively. The curve fitting results are also shown for the 200A LiF case. The experimental data is well fitted by the sum (solid line) of two peaks (dashed line), one at 685.7 eV corresponding to a LiF bonding, and the other at 688.5 eV due to C - F bond................. 148

Figure 8-3 F Is core level spectra recorded on both the (a) organic and (b) cathodesides of the cleaved interface for LiF interlayer thickness o f 15A................................... 149

Figure 8-4 XPS depth profile on the Al/LiF side o f the cleaved interface with 200A LiF layer. The evolution of the Al 2p, F Is, Li Is, O Is, and N Is core level features is shown as a function of the distance from the interface. The presence o f an Al oxide at the Al/LiF interface suggests diffusion o f O through pinholes in the Al films. Such O diffusion ends abruptly at the Al/LiF interface.................................................................. 151

Figure 8-5 F Is core level spectrum of the cleaved interface o f a single layer device a glass/SiO/Al/LiF/C6o/LiF/Al/SiO structure. Removal o f the substrate and organic layers in vacuum left behind the cathode material (SiO/Al/LiF) and approximately 50A of C6o (Binding energy values not aligned externally)..................................................155

Figure 8-6 F Is core level spectrum with high energy shoulder for (a) deposition of ~5A LiF on 350A Ceo on Si (b) deposition o f 10ML of C6o on 200A LiF on Si and(c) deposition o f 2ML C6o on ~5 A LiF on Au. The solid line in each case represents the LiF substrate, except for (a) where it represents crystalline LiF................................... 155

Figure 8-7 C Is high binding energy satellites. The dropdown lines indicate the theoretical position of the shake-up features after Enkquist et. al. [47], The oval indicates the missing p-p* feature for LiF-C6o, observed for pure deposited Cô..............162

Figure 8-8 High resolution scan of satellite structure for C(,q. Drop down lines represent the theoretically determined orbital transitions from Enkvist et al. [47] all visible in the spectrum................................................................................................................................163

Figure 8-9 Molecular energy levels of C60 (neglecting core hole ionization) (after [51 and 47]). The transitions that correspond to the observed features in the spectrum are the HOMO-LUMO transition between 5hu and 5tiu* (LUMO) at 1.9eV, and the dipole transition from 6hg to the LUMO at 6.0eV. The features at 3.8 and 4.8eV cannot be assigned to a single transition, but represent the (5hu, 7hg, and 4gg ) ->(5tiu*, 5t2U*, 8hg*, and 5gu*) monopole and dipole transitions........................................... 164

Figure 8-10 C Is shake-up structure for Ceo on (a) pure Au (b) 5A LiF coated A u...............165Figure 8-11 Evolution o f the F Is core level with C6o deposition on a variety of

substrates, (a) ~5A LiF on Ag (b) ~5A LiF on ITO (c) ~5A LiF on Au (d) 200A LiF on Si. Each cycle represents roughly 1 monolayer deposition of Cgo..........................171

Figure 8-12 C Is shake-up satellites for 2ML deposition of C6o on a variety o f substrates. 171

Figure 8-13 Normalized N Is core level for the cathode side of the cathode/organic interface for Al and LiF/Al cathodes, with TPT and Alq3 as the electron transport layer. The TPT/A1 (closed triangles) and Alqi/cathode show peaks consistent with TPT powder and the organic side of the interface, respectively..........................................173

- xv -


Figure 8-14 F Is core level for the cleaved surface, both the cathode and organic sides,for (a) Al/LiF cathodes and (b) Ag/LiF cathodes.................................................................. 174

Figure 9-1 F Is core level for LiF o f varying thickness on (a) Si (b) C6o and (c) Alq3 aligned to literature values for the underlayer. The solid vertical line at 685.6eV represents the alignment of the core level using adventitious species on the surface for all cases. Due to differential charging effects, the core levels are slightly differentfor the various substrates, but all fall within the acceptable range for ionic LiF. See text for details. Notice the broad and asymmetric peak shifted to lower binding energies for LiF on Akp surfaces attributable to charging effects......................................185

Figure 9-2 Wagner map for deposited LiF o f different thicknesses on Si, C6o and Alq3

surfaces. A majority of the points lie along lines of slope 3 indicating a similar chemical state. The difference between LiF on Alq3 and LiF on the other substrates is likely due to charging effects................................................................................................ 186

Figure 9-3 Growth of LiF on surfaces as monitored by XPS. The dotted red line is the expected change in intensity with layer by layer growth. The solid red line represents a linear sum of reduced squares best fit o f the data for thicknesses less than 100A. The intensity follows a parabolic shape (the black dashed line is just a guide to the eye), indicating initially island growth with eventual formation of a complete layer on the surfaces................................................................................................ 187

Figure 9-4 XPS Ar+ ion sputtering profiles for LiF with C6o (top row) and Alq3 (bottom row), (a) Profile of F Is core levels with a nominal LiF thickness of 3C)A on organic surfaces (b) Profile of F Is core levels with nominal thickness of 100A LiF obtained from Al/LiF cathode surface exposed by peel-off (c) Concentration profile for Al and F for the structure described in (b).................................................................................. 188

Figure 9-5 High-resolution cross-sectional SEM images of 100A LiF on (a) Alq3 and (b) C6o. To accommodate charging, the LiF on Alq3 was coated with Pt, and tilted 4° from normal.................................................................................................................................189

Figure 9-6 SEM images of the surface topography for 100A LiF on (a) Alq3 and(b) Cô surfaces. Samples tilted to 45° to image both LiF surface (top left hand side) and cross-sectional cleavage plane through Si (bottom right hand side)............................ 190

Figure 9-7 Observed charging shift as a function of thickness. Lines represent linear sum of reduced squares best fits of the data. Assuming a parallel plate model, the slope of the lines represent the electric field developed in the dielectric, given on the graph in units of MV/m.......................................................................................................192

Figure 9-8 (a) The AEF is-u is over time indicating that the chemical state is consistent with irradiation time, though LiF crystal is different from the deposited layers.(b) Change in the Li/F ratio over irradiation time indicating a slight decay due tothe formation o f F-centres The lines are just a guide to the eye...........................................196

Figure 9-9 Shift in the F Is after 45min X-ray bombardment for LiF on various substrates. The lines represent linear sum of reduced squares best fits o f the data above a critical charging shift. The cross-over point is indicated for each curve 198

- xvi -


Figure 9-10 Shift in F Is core level with irradiation time. The lines represent reduced chi2 fits of the data to a function described by equation 9-3 with Levenberg- Marquardt statistics. The right facing triangles for Alq3 indicates the change in the N Is core level with time to indicate the stability and conductive nature o f the molecule itself............................................................................................................................ 200

Figure 9-11 Change o f the F Is core level kinetic energy as a function of time. For each set of data, the first set o f lines represents the time constant derived from the non-linear curve fitting to figure 9-10. The other lines represent linear sum of reduced squares best fits o f the data for the various regions............................................... 201

Figure 9-12 Estimation of the R and C values from the linear and steady state portionsof the transient F Is core level shifts. Lines are just a guide to the eye..............................203

Figure 9-13 The calculated resistance as a function o f the charge carrier mobilities.As the capacitance o f the systems are very similar, the conductivity is related to changes in the resistance of the underlayer. The line represents a linear sum of reduced squares best fit of the electron mobility data...........................................................206

Figure 9-14 Comparison of device behaviour for 40 and 100A LiF interlayers withCgo and Alq3 based devices (adapted from [11]).................................................................. 206

Figure 10-1 Schematic o f various interface structures.............................................................. 212

Figure 10-2 Embedded oxide structure for oxidation of LiF coated Al surfaces....................215

Figure 10-3 Schematic o f interfacial structures for various metals with a LiF interlayer.... 221

Figure C-l Determination of the spectrometer resolution.........................................................239

Figure C-2 Determination of the valence band maximum (VBM)...........................................239

- xvii -


List of Tables

Table 2-1 Electrical characteristics for various cathodes.............................................................18

Table 3-1 Charging mechanism parameters (adapted from [29])............................................... 42

Table 5-1 XPS measured N/Al ratios on various buried surfaces. The sensitivity factors are: 0.472 for N Is, 0.250 for Al 2p, and 0.333 for Al 2s, respectively. The theoretical N/Al ratio is 3, calculated based on Alq3 molecular structure................................................84

Table 6-lParameters used for film structure analysis.................................................................101

Table 6-2 XPS parameters for Al 2p core level as observed on coated and uncoatedsurfaces........................................................................................................................................ 103

Table 6-3 Oxidation rates and characteristic lengths as determined by Mott-Cabreratheory (figure 6-3)......................................................................................................................107

Table 7-1 Summary o f peak positions and curve fitting parameters for coated anduncoated surfaces o f M g...........................................................................................................130

Table 7-2 Atomic ratios at the cathode side o f the as-peeled interface....................................135

Table 7-3 Summary o f peak positions and curve fitting parameters for Mg and Mg/LiFsurfaces after peel-off............................................................................................................... 136

Table 7-4 Comparison of surface lattice constants with Mg along low index planes. LiF and the products of Mg oxidation have ( lx l ) coincidence along both a and c axes.For the molecular fragments, the smallest lattice misfit is given by ( lx l ) for {1000},(2x4) for {1010}, and (3x1) for (l 102 ) planes ofMg...............................................................139

Table 8-1 XPS measured ratios on organic side o f buried surfaces..........................................146

Table 8-2 Theoretical binding energy shifts for model structure o f LiF-C6o interaction 159

Table 8-3 Gibb’s free energy of fluorination reaction at 298 K ................................................167

Table 9-1 Estimated electrical properties of the LiF film and crystal (firstapproximation - equation 9-4).................................................................................................201

Table 9-2 Estimated conductivities for LiF thin films and crystal from the transientF Is core level shift (2nd approximation).............................................................................. 204

- xviii -


Table 9-3 Dielectric properties for LiF thin films on various substrates.................................205

Table 10-1 Gibb’s free energy of metal-exchange oxidation-reduction reaction at 298 K...214

Table 10-2 Lattice constant comparisons for low index planes............................................... 218

Table E-l Theoretical bond lengths, binding energies assuming Koopman’sapproximation and Mullikan charges binding energy calculation for model structures ofLiF-C6o interaction................................................................................................................ 247

- xix -


Nomenclature

List of AcronymsAFM Atomic force microscopyAlq3 8-tris(hydroxyquinoline aluminium)ARXPS Angle resolved X-ray photoelectron spectroscopyCgo/NBB Buckminsterfullerene or NanobuckyballCMAC Central distribution chamberDFT Density functional theoryDOS Density of statesEAL Effective attenuation lengthEML Emission layerESCA Electron spectroscopy for chemical analysisETL Electron transport layerF-D Fermi-Dirac functionFEOE Full equalization of orbital energyFWHM Full width at half maximumHOMO Highest occupied molecular orbitalHREELS High resolution electron energy loss spectroscopyH-Si FIF treated Si (100) waferHTL Hole transport layerIMFP Inelastic mean free pathITO Indium tin oxideKJL Kurt J. Lesker OLED cluster toolLCD Liquid crystal displayLED Light emitting devices/diodesL-I-V Luminance-current-voltageLUMO Lowest unoccupied molecular orbitalMAC Multi-access chamberMIM Metal-insulator-metalML MonolayerMNDO Modified neglect of differential overlapMOM Metal-organic-metalNMAC Inorganic deposition chamberNPB N, M-di(naphthalene-1 -yl)-N, A'-diphenyl-benzideneOLEDs Organic light emitting diodesOMAC Organic deposition chamberPEOE Partial equalization of orbital energy

- XX -


PES Photoelectron spectroscopyPVD Physical vapour depositionSEM Scanning electron microscopySIMS Secondary ion mass spectroscopyTMFP Transport mean free pathTPD A,M-diphenyl-jV,N'-bis(3-methylphenyl) 1,1 '-biphenyl-4,4' diamine

TPP-2M Tanuma-Powell-Penn theoretical inelastic mean free pathTPT Triphenyl triazineUHV Ultrahigh vacuumUPS Ultraviolet photoelectron spectroscopyVBM Valence band maximumXPS X-ray photoelectron spectroscopya-NPD 4,4'-bis [V-l -napthyl-V-phenyl-amino]biphenyl

F-MAC Analysis chamber

List of Symbolsa ’

Cic

ae

OLMad

CCv

PX

A

AE,charging

m -j

A E m iia l

AE(t)A G °rxn

AHc

8A8 tr

S x

Modified Auger parameter

Condensation coefficient

Evaporation coefficient

Madelung constant

Vaporization coefficient

Asymmetry parameter for curve fitting

Area fraction of surface coverage

Lattice mismatch at interface

Equations where variable appears (unless indicated)3-8, 8-3

4-2

4-1

3-6

4-2

6-3, 7-1, 7-3, table 6-2

3-20, 4-3, 6-2, 6-3, 6-4

7-1, table 7-4

Shift in core level binding energy due to X-ray irradiation

Difference in core level binding energy between elements i and j

Charging shift occurring faster than time scale of XPS experiment

Transient shift in core level binding energy

Standard Gibb’s free energy of reaction

Heat of condensation

Area of an emissive source

Adatom trapping probability

Secondary electron yield from X-ray excitation

9-3, figure 9-10, figure 9-11, figure 9-13

figure 9-9(b)

9-3

9-3

table 10-1

4-5

4-6

4-33-13,3-14, 3-16,3-17, 9-4, 9-5, 9-6, table 3-1

- xxi -


e

£/£ 0

£s0

<t>B

(P

K

M71,71*

e

P ,r

<7

a, a*Os

Td

¥

Material permittivity

Dielectric constant of iVacuum permittivity in free space

Thermal emissivity of source

X-ray flux

Work function

Charge injection barrier

Emission/evaporation angle

Effective attenuation length of i photoelectron through material j

Charge mobility

Pi bond, anti-Pi bond

Electron take-off angle

Trapped charge density

Conductivity

Sigma bond, anti-bonding Sigma orbital

Stefan-Boltzmann constant

Time constant (t=RC)

Decay constant

Electron wave function (/ initial,/final)

Electric field dependent pre-factor for injection

2-1,3-9,3-10

9-7, table 3-1, table 9-3

section 9.3.4

4-6

3-13,3-14, 3-16, 3-17,9-5,9-6

3-1

2-1

4-2

3-18, 3-19, 3-20, 6-1, 6-2,6-3, 6-4, table 6-1

2-1

3-18, 3-19, 3-20, 6-1, 6-2,6-3, 6-4

table 3-1

table 9-2

4-6

3-14,3-15, 9-3, 9-4, table 9-1, table 9-2

6-8

section 3.3

2-1

at

a,1jump

cC d

Ci

Q

d

E

Lattice constant of i Ion jump distance

Geometric "capacitance" as measured by XPS

Characteristic distance for Mott-Cabrera law c-axis crystal lattice constant for i of hexagonal crystal structureRelative atomic fraction of species i

Vertical sampling depth/thickness as determined by XPS

Electric field

Fundamental electron charge

7-1, chapter 6 and 7

6-7

3-14,3-16, 9-5, table 9-2

6-5, 6-6, 6-7

chapter 7

3-2

3-18, 3-19, 3-20, 6-1, 6-2, 6-3, 6-4, 9-7, chapter 5

2-1, table 9-1

2-1,3-11,3-13,3-14,3-16,3-17, 9-1, 9-4, 9-5, 9-6

- xxii -


F°

Ed

e eff

EkEr

h

hv

h

lo

J J o

j c

J e

J in j

Jok

kj

KoxmN

Na

Nt

N0Ns

N'

p1 vap

m

m

m

Binding energy

Zero point binding energy

Desorption energy

Effective charge of defect

Kinetic energy

Relaxation energy

Layer thickness

Photon energy

Photoelectron intensity of species i

Secondary electron emission current

Current density

Condensation flux

Evaporation rate

Injection current density

Angular molecular flux from evaporation source

Boltzmann constantInteraction coefficient between core electrons and valence electrons Oxidation rate constant

Mass

Number of electronic states in atom

Avagadro’s number

Atomic density (atoms/cm3)

Density of charge hopping sites

Surface/interface atomic density

Number of metal ions per unit area available to dissolve into oxide

Equilibrium vapour pressure

Power density from condensation

Power density from kinetic energy of impinging particles

Radiation heat density

3-1, 3-3, 3-4, 3-5, 3-7,3-15, 8-1, 8-3, table 8-2

3-3

4-3

6-7

3-1,3-7,8-3

3-3,3-5,3-7,3-11,8-1,8-3

3-10, 9-7

3-1

3-2, 3-19, 3-20, 6-1,6-2, 6-3, 6-4

3-12, 3-13

6-8

4-2

4-1, 4-4, 4-5

2-1

4-2

2-1, 4-1, 4-2, 4-3, 4-4, 6-7

3-3, 3-4, 3-5

6-5, 6-6, table 6-3

4-1,4-3

section 3.3

4-5

3-19, 6-1, 6-2, 6-3, 6-4, table 6-1

2-1

4-3, table 9-1, table 9-3

section 6.3.2

4-1, 4-2, 4-3

4-5

4-4

4-6

- xxiii -


Hole charge density due to X-ray irradiation

R "Resistance" as measured by XPS

Vs(t) Surface potential (often as a function of time)

3-9,3-10,3-12, table 9-1,table 9-3

qt Charge on atom i 3-3,3-5,3-6, table 8-2Distance between source and substrate during . „ . ,deposition 4‘2' ^

3-12, 3-14,3-17, 9-4,9-6, table 9-2

ry Interatomic spacing between atoms i and j 3-6, 8-2, table 8-2

S Irradiated specimen surface area 3-12, 3-14, 3-16, 3-17

Sc Conductivity table 9-1

Si Sensitivity factor of species i 3-2

t Time 3-11,3-12,6-8

T Temperature 2 -1 ,4 -1 ,4 -2 ,4 -3 ,4 -4 ,4 -6 ,6 -7

V Potential 3-9

Vc Coulomb potential 3-3

Vcontact Contact potential 6-7

Madelung potential 3'3 ,3 '6’ 3-11»8-1»8'2’table 8-3

3-10, 3-11,3-12, 3-14,3-15, 3-16, 3-17, 9-4, 9-5, 9-6

x Oxide thickness for Mott-Cabrera law 6-5, 6-6

z Distance between analyzer and substrate during XPS 3-9

Z°- Empirical zero-point energy 3-4

§ Change in oxide thickness over time 6-5

- xxiv -


Chapter 1

Introduction

1.1 Interfaces in organic electronics

Although organic molecules are traditionally thought o f as insulators, organic conduction has

been intensely studied over the last 50 years to capitalize on the photoconductivity o f many

molecules under visible light [1], Manufacturing devices using organic molecules cheaply,

under ambient conditions, with a high degree of property flexibility, would represent a

fundamental, yet desired change for microelectronics. However, the performance o f organic

devices has lagged behind those based on traditional semiconductors due to instability and

low purity o f many organic semiconductors; intrinsically low carrier mobilities and difficulty

of doping, and difficulties in making reliable electrical contacts [1], Within the realm of

optoelectronics, however, organic semiconductors can be considered a viable alternative as

the performance has become comparable to or even surpassed more traditional materials [2].

- 1 -


Chapter 1 Introduction 2

Electroluminescence from organic molecules, first observed by Bemanose in the

1950s in crystalline thin films of acridine orange and quinacrine under a high-voltage AC

field [3], has long been of interest due to the possibility o f high fluorescent yields throughout

the visible range [4], Early attempts at electroluminescent devices by Pope et al. [5], and

Helfrich and Schneider [6], however, showed poor charge injection into the organic single

crystals [7] and high operating voltages due to impurities, which prohibited most commercial

applications [3,7]. The discoveries o f efficient electroluminescent organic devices based on

thermally evaporated small organic molecules [8,9] and conjugated polymers [10] led to a

major revival of interest in organic light emitting diodes (OLEDs) [1].

The primary motivation for the development of OLEDs has been their potential as the

next generation display technology, due to their high brightness, high viewing angle, full

spectrum colour, thinness, and low driving voltage [11], High quality commercial products,

such as a 40” television from Samsung [12], have already been produced. The flat panel

display market in 2006 is estimated at US$57 billion, with OLED based displays making up

4% o f the current market. The desire to improve that market share has provided the incentive

to continue optimizing device performance [13], though the growth of OLED technologies

has been mainly limited to Asian markets. While interest in commercialization has declined

in recent years due to the success of LCD technologies in large area flat panel displays, the

improvements in digital broadcasting, and the continued miniaturization of displays for

handheld consumer electronics, have reinvigorated the research in OLEDs. The future

success o f OLEDs, however, rests on their potential for use on flexible substrates. The

driving force for continued research, especially in the characteristics of the organic/electrode

interfaces in OLEDs, currently rests on the desire to make fully flexible displays [1], which

are difficult to fabricate with other display technologies.



Within the OLED field, the research emphasis to date has mainly been on the

development of highly efficient devices over a range o f colours, through modifications to the

emission properties of the organic molecules. Throughout the last twenty years, a wide

variety o f organic molecules and polymers have been used as the active emitting layers [14]

in an attempt to produce the entire visible spectrum. In order to ensure adequate device

performance, an even wider variety o f inorganic materials have been attempted as electrical

contacts with these different molecules. These studies have established that the interfacial

region between the organic active layers and the inorganic contacts plays a primary role in

device performance, through the control o f effective carrier injection and long term device

reliability. Since injection and reliability are two key factors in the efficiency o f devices, the

contact formation at the interface represents a critical feature in the widespread commercial

application of OLEDs. However, unlike inorganic semiconductor/metal systems, where

contact formation has been studied extensively [15], organic/inorganic interfaces in these

systems are not fully understood.

Therefore, a clear picture o f the chemical, physical, and electronic structures at the

organic interface with inorganic contacts would be helpful in understanding device behaviour

and optimizing device performance. Studies on the organic/inorganic interface in OLEDs for

a variety of systems have been undertaken by others. These studies have indicated that a wide

range of interfacial types are possible in OLEDs. The relation between the interfacial

structure and the device performance, however, is still the subject of some controversy and a

number of conflicting mechanisms have been proposed. One of the major limitations o f these

previous interfacial studies has been the focus on ideal monolayer structures. The most

commonly used techniques have limited ability to examine the buried interfacial structures

that occur in conventionally fabricated devices. One of the aims of the present work,



therefore, is to provide a connection between the traditional approaches to interface

characterization and the manufacturing conditions used for typical OLED structures. The

desire is to compare the interfacial structure in devices, analysed by the novel application o f a

traditional adhesive tape test [16], with the structures developed during monolayer deposition

to gain a deeper understanding of the interface formation process. Idealized interfaces as

observed by studies with monolayer growth may not be as relevant to those developed under

real manufacturing conditions that ultimately affect device performance. In these studies, the

predominant tool for analysis was X-ray photoelectron spectroscopy (XPS).

1.2 Thesis organization

After a brief overview o f organic light emitting diodes and the role o f the cathode/organic

interface in the next chapter, this thesis continues with an introduction to X-ray photoelectron

spectroscopy, and its previous use for studies of the buried interface in OLEDs in chapter

three. Following is a description of the experimental methods and instrumentation used in

this project. The remaining chapters deal with the experimentation carried out in this project

to describe the interfacial chemical structure, beginning with simple metal/Alq3 interfaces in

chapter five. With the introduction of a LiF interlayer, however, a simple metal/organic

junction description is no longer sufficient to portray the interfacial structure. Rather, the

cathode/organic junction is complicated by the existence o f two interfaces - that between the

metal and the interlayer, and that between the interlayer and the organic. The rest o f this

thesis, therefore, examines these interfaces and the interlayer itself in turn. Chapter six and

seven deal with the impact of LiF on the metal surface, focussing on the oxidation of A1 and

Mg, respectively. The formation of charge transfer compounds between organic molecules



and LiF is introduced in the first part of chapter 8 , and described in greater detail for the

specific case of Ceo-LiF interaction in the second part. As is discussed in chapter 9, the

properties o f the LiF layer itself as the thickness increases cannot be described in isolation

from the other components. From the results o f these examinations, models of the interfacial

structure in organic devices are proposed and the role o f LiF is described in chapter ten.

1.3 References

1 S. Forrest, P. Burrows, and M. Thompson, IEEE Spectrum Aug, 29 (2000).

2 J.R. Sheats, H. Antoniadis, M. Hueschen, W Leonard, J. Miller, R. Moon, D. Roitman, and A. Stocking, Science 273, 884 (1996).

3 M. T. Bemius, M. Inbasekaran, J. O’Brien, and W. Wu, Adv. Mater. 12, 1737 (2000).

4 Y. Sato, in Electroluminescence I , edited by Gerd Meuller (Academic Press, San Deigo, 1999), Vol. 64 Semiconductors and Semimetals, Chap. 4, p.209.

5 M. Pope, H. P. Kallmann, and P. Magnante, J Chem. Phys. 38, 2042 (1963).

6 W. Helfrich, W. and W.G. Schneider. Phys. Rev. Lett. 14, 229 (1965).

7 M. D’lorio. Can. J. Phys. 78, 231 (2000).

8 C. W. Tang and S.A. VanSlyke, Appl. Phys. Lett. 51, 913 (1987).

9 C.W. Tang, S.A. vanSlyke, and C.H. Chen, J. Appl. Phys. 65, 3610 (1989).

10 J.H. Burroughes, D.D.C Bradley, A.R. Brown, R.N. Marks, K. Mackay, R.H. Friend, P.L Bums, and A.B. Holmes, Nature 347, 539 (1990).

11 For recent reviews of OLEDs and OLED prospects see P. Burrows, S. R. Forrest, and M. E. Thompson, Curr. Opin. Solid State Mater. Sci. 2, 236 (1997); (b.) ref [1] above; (c.) L.J. Rothberg, and A.J Lovinger, J. Mater. Res 11, 3174 (1996); (d.) L. S. Hung and C. H. Chen, Mater. Sci. Eng., R 39, 143 (2002).

12Samsung (May 19, 2005) “SAMSUNG Electronics Develops World’s First 40-inch a-Si-based OLED for Ultra-slim, Ultra-sharp Large TVs” Press release.

13H. Antoniadis, Thin Film Users Group Proceedings, Sunnyvale, CA, (May 2003). American Vacuum Society - N. California Chapter.

14 U. Mitschke, and P. Bauerle, J. Mater. Chem. 10, 1471 (2000).

15 see for example Metal-Semiconductor Schottky Barrier Junctions and Their Applications, edited by B. L. Sharama (Plenum, New York, 1984).

16 M. Ohring, The Materials Science o f Thin Films (Academic, Toronto, 1992), p. 444.


Chapter 2

OLED fundamentals

2.1 OLEDs and organic conductors

In contrast to the mature inorganic LED field, organic light emitting devices (OLEDs) are

just recently beginning to show real promise, as display applications are finally being

realized. Light emitting devices based on organic molecules are referred to as light emitting

diodes due to the non-linear current rectification, with a minimum tum-on voltage, of the

organic materials, which are analogous to the forward voltage drop o f an inorganic diode [1].

2.1.1 Device operation

Organic light emitting materials were first proposed as an alternative to inorganics in LEDs

due to their versatility, both in range of emission and complexity o f device patterning on a

wide variety of substrates, and in the potential ease o f processing under ambient conditions

- 6 -


Chapter 2 OLED Fundamentals 7

[2], OLED operation is mainly controlled by the injection characteristics, due both to the

nature of organic conduction and luminescence, and to the limited thickness of typical

devices.

2.1.1.1 Conduction

Conduction in organic molecules can be characterized as thermally activated three

dimensional variable range [3] hopping mechanism through conjugated 7t-bonds, due to the

localization of electronic states to the individual molecules [4], As a consequence of the

relative disorder o f molecular materials, and the localization o f electronic states, organic

semiconductors lack intrinsic charge carriers that can contribute to conduction and

luminescence. The introduction of charge carriers, therefore, depends solely on the injection

characteristics o f the electrode contact. Baldo et al. [4] showed that the broad distribution of

hopping sites at the electrode interface creates sufficient disorder in the organic layer to

dominate the transport characteristics due to the limited thickness o f the transport layers.

2.1.1.2 Luminance

Unlike in traditional semiconductors, luminance in OLEDs is achieved through the de

excitation or recombination o f a bound electron-hole pair (exciton). The organic molecules

act as charge carrier traps, which then attract the oppositely charged carrier to form an

exciton [5]. Light is then emitted from exciton decay as if from an excited molecule, with

relaxation from the excited state of the molecule, known as the lowest unoccupied molecular

orbital (LUMO), to the ground state, the highest occupied molecular orbital (HOMO).

Similar to conduction, luminance is explicitly linked to the injection properties, since the size

and position of the recombination zone are affected by rate of electron injection [6 ], The

work function of the electrode controls the size of the internal electric field, which in turn



controls the mobility o f charge carriers, and hence the position of the recombination zone. As

well, since the electrode can act to quench excited molecules [7], the device efficiency

increases with the distance the charge carriers potentially travel before combining to form an

exciton. Since luminescence is a result o f exciton decay, requiring electrons and holes, the

injection properties at both cathode and anode interfaces play a major role in device

performance. However, if the number o f carriers is not balanced, the recombination and

device efficiencies are limited by the number o f minority charge carriers [8 ], As many

organic molecules are predominantly hole transporting and easily doped with holes, the

injection o f electrons at the cathode/organic interface is the limiting factor in device

efficiency and driving voltage [9].

2.1.2 Device structures

The basic OLED, shown in figure 2-1, consists o f a multilayer structure built up on a glass

substrate, with the active organic films sandwiched between two dissimilar electrodes [1 ,1 0 ].

For typical device configurations, the anode is deposited first, directly onto the substrate.

Anodes tend to be thin transparent conductive films o f high work function suitable for hole

injection (typically indium tin oxide (ITO)). Atop the anode are deposited one or more

organic layers, each between 500 to 1000A, that can act as hole-transporting, emitting, and

electron transporting layers. The cathode is then evaporated on top of these organic layers.

Multiple evaporation steps may be required for cathode formation if multilayer cathodes are

desired. Finally, the device is usually encapsulated to prevent oxidation o f the various layers.

The thickness of these devices, generally between 0.1 to 2.2/um, magnifies the effect o f the


Chapter 2 OLED Fundamentals

interface on the device properties, since the interfacial interaction region can dominate the

thickness of the active layers.

cathodest

2-10 VDC

substrate (glass)

, electron transport layer (ETL)

• - light emitting layer (EM L)

hole transport layer (HTL)

anode

light

Figure 2-1 Schematic of typical OLED structure

2.1.3 Organic electron conduction layers:

2.1.3.1 8 -tris(hydroxyquinoline aluminum) (Alqs)

Introduced in 1987 [11] in the first stable OLED device produced, 8 -tris (hydroxyquinoline

aluminum) (Alqs) has remained the prototypical and most widely used small organic

molecule for OLED applications. Belonging to a class of metal chelates [12], ALp is a

symmetric organometallic molecule with an A1 ion surrounded by three 8 -hydroxyquinoline

ligands (figure 2-2). Al, therefore, is in a 3+ oxidation state, having electrons localized on the

O on the phenoxide ring of the ligand, with relatively weak bonding to N in the pyridinal ring

[12]. Produced by thermal evaporation, Alq3 tends to form smooth, pin-hole-defect free thin

films with crystalline domains smaller than 500A [13]. Though films are a mixture o f the two

geometric isomers, the observed structure is predominantly meridinal [14], with the three

quinolate rings perpendicular to each other, rather than in-plane. Films of thicknesses typical

for OLEDs show no diffraction peaks; however, the molecular packing is preferentially that

o f triclinic a-Alq3 crystals, suggesting an intermolecular spacing o f approximately 20A [15].



V ^ s . -

s. > “N.-if 4 l

(b)

Figure 2-2 Alq3 molecule (a) planar structure (b) Three-dimensional model of the meridinal isomer [12]

Relatively unique among organic molecules, Alq3 tends to be a strong electron

< >ytransporting material, showing field dependent mobilities as high as 1 0 ' cm /V s at

intermediate electric fields expected for devices1 [16]. Due to its high mobility, Alq3 is often

used in devices as a combination electron transporting and emitting layer. Alq3 emits most

efficiently in the green region, around 550nm [10]. However, the emission characteristics can

be tuned by doping [13], by modification o f the ligands [2 ] or by coordination around the

central A1 ion [2]. Alq3 is one of the most fluorescent and stable molecules in the class of

metal chelates [13], but its relatively low fluorescence yield (8 %) results in an upper limit of

2% electroluminescence efficiency in lumens/watt (lm/W) [11]. Though this yield is

comparable to commercially available light emitting diodes, it is still inferior to standard

light sources. It has been difficult to find another organic molecule that surpasses Alq3 in

electroluminescence, electron transport [16] and stability properties [17]. Therefore, Alq3

remains the most widely used small molecular weight organic molecule, representing the

archetypal material for studies of organic electron transport [4] and device performance. As

such, it can be considered analogous to Si in traditional inorganic semiconductor devices.

1 2 Though relatively low compared to inorganic materials such as Si (p = 1400 cm /Vs,) or even amorphous Si-3 7(p= 10 cm /Vs), this electron mobility is high for organic materials, which tend to be hole transporting.



2.1.3.2 Nanobuckyball (Cso or NBB)

Cgo is an allotrope of carbon consisting 60 C atoms arrayed in a truncated icosahedron, with

12 pentagons and 20 hexagons, as shown in figure 2-3. C6o is one o f the highest symmetry

molecules, with 120 operations for icosahedral symmetry, and can be considered as almost

spherical [18]. Consisting of a K-conjugated network, C6o has interesting electrical and

chemical properties, even though it is relatively non-aromatic [19]. Less thermodynamically

stable than either diamond or graphite [18], C60 is the purest source o f carbon, without any

attached functional groups or dangling bonds for interaction with the surroundings. This

makes it the ideal molecule to study the potential for bond formation between C and

inorganic materials, and to examine cathode metal oxidation from oxygen sources other than

the organic molecule itself.

Figure 2-3 C60 molecule

Once C6o powder is produced by anaerobic combustion or pyrolysis o f aromatic

hydrocarbons, or by arc-discharge [18], the cage structure of the molecules is fairly robust,

and high quality films can be produced by thermal evaporation. According to Halac et al.

[20], for deposition energy less than 20eV, the molecules would have very little cage

distortion upon deposition, as in the Monte Carlo simulations of figure 2-4 (a) and (b).

During thermal evaporation, with an average deposition rate of 1 A/s and deposition source

temperatures of ~400°C, the average impact energy for the C6o molecules is 0.1 eV.


Chapter 2 OLED Fundamentals

(a) (b)

12

1 ev

300 ev

Figure 2-4 Monte Carlo simulations of C6o growth [20] on (a) Si and (b) diamond substrates (c) low density films grown by throwing C6o molecules at Si substrates at low energy

300 eV20 eV

Although the molecules maintain their shape, dense film formation is not always

guaranteed during evaporation. Dependent on the substrate used, below an impact energy o f

50eV, the film may form a porous structure with large intermolecular holes [20], as seen in

the simulation o f figure 2-4(c) above. The film formation process, however, is highly

dependent on the substrate used, and fully dense epitaxial films are possible [21]. Regardless

of the film density, C6o films formed by thermal evaporation do replicate the structure o f the

fulleride crystal. X-ray diffraction measurements, as in figure 2-5, indicate sharp FCC

diffraction lines, without a significant amorphous fraction, as seen by the low background in

the spectra. However, as expected there is no long range order in the crystalline domains.

Hebard et al. estimate the coherence length of C6o to be about 60A or 4 unit cells [22], The

structure of thermally evaporated C6o films, therefore, consists of relatively intact molecules

that pack as spheres, partially replicating the FCC structure of fulleride crystals [22],

*r

3 1000

Cm (4450 A)

222

5 500331

400 333

25_j_30

Figure 2-5 X-ray scan of a 4450A thick C7,o film deposited onto an off- axis cut single-crystal sapphire substrate. The markers indicate the calculated diffraction lines from a face-centred-cubic cell seen in diffraction from bulk C60 powder [22].

35



Since its discovery two decades ago [23], the electrical properties of C6o thin films

have been intensely investigated. The variable nature of C6o, effectively acting as both an

electron and hole transporting semiconductor, makes it suitable for a wide range o f

applications. Due to high carrier mobilities for both electron and holes, 1.3 and 2x1(L4

cm2/Vs respectively [24], comparable to those of amorphous silicon, C(,q thin films have been

utilized as the active element in photovoltaic solar cells [25], thin-film field effect transistors

[26], rectifying diodes, [27] and more recently as an electron transport layer for organic light

emitting devices [28], Stable cathodic metals, such as A1 and Al/LiF, have been used very

effectively in such devices, with quasi-ohmic behaviour in the case o f Al/LiF [28], due to the

low electron-injection barriers at the interface for C6o electron-transport layers. However,

such devices are extremely susceptible to oxidation. The contact degrades from ohmic to

blocking after exposure to air due, to the emergence of a potential barrier at the interface

[29]. Moreover, a reduction in conductivity of several orders o f magnitudes due to oxygen

adsorption has been reported [30,31]. To achieve reliable and robust devices, encapsulation is

usually required to prevent degradation of C6o devices due to oxygen exposure [32],

Nevertheless, due to its ability to act as both an electron acceptor and donor, it is becoming a

standard material as both a dopant and interlayer for organic optoelectronics.

2.2 Role of the interface in OLEDs

The interface plays a central role in understanding device performance, both during initial

operation and over time. Its importance can be related primarily to the major effect of the

interfacial properties on carrier injection, and therefore as described above on effective

conduction and luminescence, and on long term device reliability.



2.2.1 Injection

Due to the observed temperature dependence o f the current-voltage characteristics of devices,

electron injection into the organic layer is presumed to be controlled mainly by a thermionic

emission process, governed by a modified Richardson equation, as follows [33]

J i n j = J0eilE expkT

exp e*E\ 2

Vv4 n e {k T )

(2 -1)

where £ is a function of the electric field, N 0 is the density o f charge hopping sites, is the barrier to injection, E is the applied electric field strength, ji is the electron mobility, e is the material permittivity and e is the electron charge.

Most of the variables controlling this mechanism are set either by the device

conditions, such as the applied electric field and operating temperature, or by the organic

layer properties, such as the dielectric constant or thickness or mobility. Therefore, it is the

interfacial conditions - the barrier to charge injection and the density of interfacial sites —

that control the modification o f injection properties. Baldo et al. [4] have indicated that the

barrier to charge injection can be related to the formation o f a dipole at the interface, with

injection from this dipole region into the organic being the limiting mechanism. This is

supported by experimental evidence of band bending, showing that the HOMO-metal Fermi

energy band offset was independent o f the metal work function for Alq3 [34] and for Cgo

[35], Other effects such as interfacial defect states and tunnelling barriers, which have also

been proposed as possible injection mechanisms [36,37,38], are also confined to the

interfacial region.

Figure 2-6 summarizes the various interactions between surfaces that may modify the

electronic charge balance and result in dipole formation. Since these can also be affected by

the nature o f the preparation of the films themselves [39], the overall quality of the interface

also becomes o f interest. Therefore, it is important to characterize the electronic structure, the



chemical state, and the morphology o f deposited polycrystalline metals and amorphous or

polycrystalline organic layers to understand the effect of the interface formation process on

the device properties [40].

Cation

I

(ah

AnionFormation Formation

;an- MP*

m

Mirror SurfaceForce Rearrangement

(b) W

ChemicalInteraction

InterfaceState

PermanentDipoleWj

Cd> (9)

®)

<0

Figure 2-6 Various substrate/molecule interactions that can lead to dipole formation [39]

2.2.2 Device reliability

Beyond the effect that charge injection has on device performance, the cathode/organic

interface plays a crucial role in long term device reliability. Organic devices are beginning to

show lifetimes comparable to those o f the conventional emissive media used for display

purposes [41]. Unfortunately, the lack of long-term stability is still one o f the major barriers

to their widespread commercialization. Device degradation falls into three broad categories

[42], The first two involve the thermal or electrochemical breakdown of the organic layers

through crystallization [43], hole induced damage in electron transporting layers [44], or

interdiffusion [45]. The third involves degradation that occurs at or as a result o f interfacial

contacts [46]. The possibility of chemical reactions that can take place at the interface, such

as the destructive reaction of the organic layer induced by the electrode [47], or the

electrochemical coupling of cathode alloy components [48], can greatly affect both the long



and short term device characteristics. As well, the surface morphology and interfacial defects

can act to limit device performance through cathode delamination [47,48,49,50,51,52], or

pinhole formation [49,51], which are often linked to interfacial reactions. Though it is the

environmental, thermal and electrical degradation o f the organic layers that have the most

obvious effect on the luminance of the device over time, the interfacial effects play a critical

role by accelerating or inhibiting the detrimental effects.

2.3 Cathode performance

2.3.1 Elemental metal cathodes

Historically, investigations into the effect of various cathodes on device performance were

approached in the context o f accepted models o f inorganic semiconductor interfaces [53],

especially regarding the injection processes. With little understanding of the nature of

organic/metal interfaces in these systems, the barrier to charge injection in the injection

process was presumed to simply be the difference between the LUMO of the organic

molecule and the work function of the metal cathode. There were attempts to utilize a

number of low work function metals based on a work function matching scheme. However,

relatively high work function cathodes such as Mg showed high efficiency and good device

performance compared to lower work function cathodes [54,55]. As well, below a threshold

value, the efficiency was in fact seen to decrease slightly with decreasing work function, as

seen in figure 2-7 [54], Since Mg is thought to form an ohmic contact with Alq3 [56], the

criterion of work function difference is inadequate to fully describe the barrier height

controlling the injection properties.



Figure 2-7 Device performance as a function of the cathode metal work function (a) relative luminance at a constant current density (b) relative efficiency at a constant luminance [54]

3.0 3.5 4.0 4.5

Work function (eV)

These low work function cathodes, due to the ease of electron stripping necessary for

high injection efficiency, are unfortunately also highly unstable, and are therefore prone to

oxidative or corrosive attack by the organic layers or by atmospheric gases. To mitigate these

effects, attempts were made to modify the cathode materials, by alloying with more stable

metals or by introducing interfacial buffer layers.

2.3.2 Alloy cathodes

Atmospheric oxidation tends to decrease the lifetime of devices, through degradation of the

cathode into an insulating oxide. To improve the stability, low work function cathodes were

coupled with more stable, higher work function metals such as A1 or Ag. Small amounts of

alloying constituents were found improve the device lifetime by orders o f magnitude [11,

55,57]. The first viable device lasting more than a few hours, shown by Tang and Van Slyke

in 1987, was in fact a bi-layer device with an Alq3 emitting layer and a Mg:Ag alloy cathode

[11]. Table 2-1 lists the relative performance of a few metal cathodes, showing the

Yb<Do £ 0.6 cE 0.4 3 —I

Sm

Zn

0.2

Cu0.0

> . 'o0 0.6 O

i5 0.4Zn

0.2



superiority of alloy cathodes over elemental ones. Though lifetime measurements are not

shown, both alloys show orders o f magnitude longer lifetimes than the pure metal Mg and Li

cathodes, which degraded in a matter o f seconds or hours [55, 58], Kim et al. [58] found that

there was also major improvement in the tum-on voltage, with very little change in the work

function for the elemental and alloyed cathodes.

However, as these alloy cathodes are produced by co-evaporation of the two metals,

reproducibility o f the correct ratio for optimal device performance is challenging to achieve

[59], leading to uneven performance for devices and subsequently to low device yields. The

device performance for the Mg:Ag cathode taken from Aziz et al. [44] shown in table 2-1

likely represents the mid-range of values for these cathodes, with poorer performance than

for the optimized Mg cathode devices produced by our group. Nonetheless, the Mg:Ag

couple first used by Tang and Van Slyke and the Li:Al couple [60] are considered to have the

best device performance for purely metallic cathodes, and as such, are still widely used.

Table 2-1 Electrical characteristics for various cathodes2

Al Au Mg Li* Mg:Ag# Li:Al LiF/Ale MgFj/Af4 Li20/Al***

Tum-on voltage (V) 4.5 7f 2.9 6 3 3.5* 2.6 4.3 3.5

Luminance At 7V (Cd/m2) 50 l f 6000 4 670 400* 9000 300 2000

Current density at 7V (A/m2) 40 - 3000 - 185 1080** 4000 75 170

Max current efficiency

(Cd/A) (at V)

1.75(10V) 0.0165* 2.4

(6V) -3

(8.5V) - 3.0(5.5V)

4.4(7V)

5.8 (6.5 V)

1 Estimated from Kwon et al. [61] t Estimated from Mason et al. [59] "Estimated from Aziz et al. [44] *Estimated from Haskal et al. [55] ** Estimated from Kim et al. [58] ***Estimated from W akimoto et al. [62] Êstimated from Fujikawa et al. [63], eLiF nominal thickness 5A

2It is somewhat misleading to compare devices produced by different research groups directly. Our group has

produced some of the most efficient and brightest devices with the lowest tum-on voltages, even for relatively poor cathodes like Al. Typical literature values range around 6.9V [59], but can go as high as 14V [37].



2.3.3 Bi-layer cathodes

In addition to susceptibility to atmospheric oxidation, low work function metals may also

react with the organic layer itself. As well, due to the small size and relatively high thermal

energy of the evaporated metal, the impinging “hot” atoms can diffuse into the weakly

bonded organic layers [64]. This can severely limit their injection behaviour, such as the case

of photoluminescence quenching phenomenon observed with metal diffusion from Ca

cathodes in oligomers [65]. The introduction of interfacial oxide layers, such as CaO for Ca

[6 6 ] and AI2O3 for Al [67], has led to significant improvements in the performance for those

cathodes. This improvement was attributed to the prevention of uncontrolled reactions at the

interface. Recent work by Kiy et al. showing improved efficiency for Mg cathodes operated

in air compared to those devices grown and analyzed in vacuum where there is no possibility

of cathode oxidation [57] indicate that interfacial oxides may in fact be required for optimal

device performance.

The superior performance of devices using ultrathin layers o f LiF or MgO with Al

[36] also suggests the importance of these interfacial insulators. The shift to multilayer

cathodes has opened up the possibility of utilizing more stable metals, such as Al and Ag. Al,

showing poor injection characteristics by itself, was of especial interest due to its inherent

stability from the formation of a passivating oxide film and high compatibility with Si

integrated circuit technology used for displays [36].

The role that these interlayers play though widely studied, has still not been

completely established. Beyond the prevention o f interfacial reactions, there have been a

number of interpretations brought forth to explain the mechanism behind the improved

injection, including electron tunnelling through a thin insulator [36], band bending at the

cathode/organic interface [36], lowering of the cathode metal work function [6 8 ], introducing



interfacial dipoles [4], or doping o f the organic layer with ions dissociated from the

interfacial compound [59]. Using these assumed mechanisms as a guide, researchers

attempted a number o f other compounds with an Al overlayer in particular, ranging from

alkali and alkali earth fluorides [59,62,69] to doped organic layers [70] and organometallic

molecules [38,70], some of which are shown in table 2-1 above. The original LiF/Al bilayer

introduced by Hung et al. [36] has proved to be the best cathode for Alq3 based devices due

to its superior properties and ease o f reproducible fabrication compared to alloy cathodes.

Initially it was also assumed that the lowered work function at the interface, if it

exists, may be attributed solely to the interlayer itself. Therefore, there was some interest in

using the stable oxides alone as cathode materials to eliminate the need for two deposition

steps in manufacturing [38,59] by exploiting the fact that the work function of a compound

can be considered to depend mainly on the work function o f the element o f lower

electronegativity [71]. However, the various oxides and fluorides attempted could only be

used as ultrathin layers [6,36,38,53,59,61,67]. Optimal thickness observed for LiF (5A) [6,

36,61], alkali and alkali-metal acetates (5A) [38], alkali metal compounds (3-10A) [59], and

AI2O3 (12A) [67] indicate that interlayer thickness is generally limited to <10A for effective

device performance. At such a thickness, the insulator is inadequate to protect the organic

from oxidative attack by the ambient environment. It was assumed, therefore, that one o f the

main purposes of the metal capping layer was to provide protection against environmental

degradation of the organic layer [38,61]. However, subsequent studies have shown that

certain bi-layer combinations, such as Mg/LiF (figure 2-8(a)) [9,72] and Al/Ge02 (figure 2-

8(b)) [36], have a negative effect on device performance compared to pure elemental

cathodes, or others, such as MgiAg/AloCh [73] and Ag/LiF [37], have no effect at all.



— OAL i F / Mg — □— 10A Li F/ Mg — OA LiF/AI — o — 10A L iF /A I

0 .4

0 .3

0 .2

0.1

0 .00 2 4 6 8 10 12 14 16

V o lta g e (V)

1.2

1.0

C a0.8

bm

0.4

Zna w ith o u t LiF o w ith 1 nm LiF0.2

0.02.5 3.0 3.5 4.0 4.5

W ork function (eV )

1000

AI/MgO Mg09Ag01 A|

./ /— 100

AI/GeO,

5 10 15 20

Drive vo ltage (V)

Figure 2-8 Effect of interlayer with a variety of cathodes (a) for current density and (b) for relative efficiency for a LiF interlayer from Stofiel et a l [9], (c) MgO and G e02 interlayers from Hung et a l [36].

The many investigations into the impact o f an interlayer between the organic and the

metal cathode have resulted in the proposal o f many conflicting mechanisms, for all o f which

there is both supporting and contradicting evidence. Although Al/LiF has proved to be the

best cathode combination for Alq3 based devices, other bi-layer combinations, such as

CsF/Al and LiF/Ca/Al, are more effective with polymeric conducting layers [74,75], It

appears that the interactions of the metal capping layer, the interlayer, and the organic layer

could be the controlling factor in charge injection. Although there have been a number of

cathodes attempted and a number of possible interpretations suggested, there is still little

understanding of the physical reasons for the superiority of certain cathode combinations.

The effectiveness o f bilayer cathodes with relatively thin interlayers indicates that the

interfacial region is indeed of primary importance for controlling device performance.



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Chapter 3

Theoretical background of X-ray photoelectron spectroscopy and its applicability to interfacial analysis in OLEDs

3.1 Basic principles

The single most productive experimental method to study both the chemical and electronic

structures o f organic molecules and the interfaces they form with the cathode materials has

been photoelectron spectroscopy [1]. Photoelectron spectroscopy (PES) is based on the

photoelectric response of a surface to the bombardment with an electromagnetic energy

source. In photoemission (figure 3-1), a photon o f energy hv penetrates the surface and

optically excites an electron from an initial state to a final state. The excited state then

propagates through the material to the surface and is ejected into the vacuum.

-25 -


Chapter 3 X-ray photoelectron spectroscopy 26

9 Photo electron

ValenceBand

Photon

Core hole<!>

Core levels

-vac

- E p —

- O

tic Energy

Energy

Figure 3-1 Photoemission process

The binding energy can be estimated from the measured kinetic energy o f the ejected

electron from

Eb = h v - E k -(j) p .!)

where Eb is the binding energy of the atomic orbital from which the electron originates, h v is the energy of the incoming photon, Ek is the kinetic energy of the ejected photoelectron and (j) is the spectrometer work function ( EVAC - EFermj).

If the energy distribution of the emitted electron is plotted against the estimated

binding energy from equation 3-1 above, the number of emitted electrons per energy interval

gives a replica o f the electron energy distribution in the solid surface [1,2], shown in figure

3-2. Depending on the energy of the photon source, either valence band or core level

electrons may be excited, as in figure 3-1 above. By using a low energy excitation source (5-

100 eV), such as an ultraviolet light (UPS), the improved resolution allows examination of

the fine electronic structure in the valence band.



Euin Spectrum

Valence Band

Sample

- V a c u u m / u e v e l - - N(E)

hu

Core Level

Figure 3-2 Relation between the energy levels in a solid and the electron-energy distribution in the photoemitted spectrum. The excitation source determines the range of interest [2].

A higher excitation energy source (>1000eV), such as X-rays (XPS), enables identification of

the chemical constituents in any given sample, since the core level electron configuration is

unique for every element (see figure 3-3). For this reason, XPS was historically referred to

as Electron Spectroscopy for Chemical Analysis (ESCA) [3].

Q ) 4 U U U COoo25 1000c .3o ( >

0

1s BINDING ENERGIES

Li Be B

*4 *4 i1/2 ■1/3

J____I___ L.

'1/3

_J I I l_200 400 6 00

B in d in g E n e r g y ( e V )

Figure 3-3 Core levels for various elements [2]



The relative intensities of the core levels for different elements can also be used to

determine the relative atomic fraction of any element, Cx, by

where I is the peak intensity and S is the sensitivity factor.

Therefore, PES techniques, particularly XPS, are especially well suited to the study o f

thin films of organic molecules and interfaces, since they provide the maximum amount of

information on the chemical bonding structure as well as some information on the electronic

valence structure. In addition, they are generally less destructive to organic systems than

electron spectroscopies and are surface sensitive, due to the short attenuation length of

escaping electrons.

3.2 Chemical shift

3.2.1 Prediction o f chemical shift from absolute binding energy calculations

The position of the orbitals in an atom is very sensitive to changes in the chemical

environment of that atom. As the overall charge on the atom changes, the energy o f the

remaining electrons is also changed. After the pioneering work of Siegbahn on XPS in the

early 1950’s [3], it has been widely acknowledged that there is a relationship between the

chemical environment, as represented by the atomic charge, and the binding energy as

determined by XPS.

During photoionization, the redistribution of energy with a change in charge

configuration from that of the neutral atom appears in the spectrum as shifts in the measured



position of the binding energy. The shift to higher binding energy with functional group

addition to a polymer background, for example, is easily understood in a simple initial-state

picture: the highly electronegative atoms in the functional groups withdraw electrons and this

reduces the charge density particularly on the covalently bound atoms of the functional

group. The core level binding energies are therefore increased. Additionally, there might be a

difference between other species in the dynamic response of the multi-electron system to the

creation of a core hole (screening effect). Both of these effects are usually referred to as the

“chemical shift” [4]. The strength of XPS as a characterization technique rests on this ability

to predict the chemical environment o f the constituent elements. If there is an established

standard for the binding energy o f the neutral atom, the observed value o f the binding energy

can be used to predict the charge distribution of the probed element.

One of the most useful descriptions of this relationship is the classical charge

potential theory of Siegbahn et al.[3]. In this theory, the binding energy of the core level can

be expressed by

where Ey is the binding energy as determined by XPS, E b° is the zero-point energy, qj is the

quantum chemically calculated atomic charge, kj is the interaction coefficient between the core level and valence level electrons, Vc is the Coulomb potential contributed by all the other atoms in the molecule due to their charge, and ER is the relaxation energy due to screening.

From a practical perspective, empirical linear relationships for a number of elements have

been developed in the literature (see the chart in Appendix A for various relationships):

where ZJ is the empirically determined “zero-point energy” from the intercept of a linear regression analysis.

(3-3)

(3-4)



For these empirical equations, the initial state effects are considered to be the major

contributing factor to the chemical shift. The “zero-point” intercept, therefore, incorporates

the effects of the surroundings and relaxation without explicitly defining them individually

[5], Nevertheless, the equations may be used to correlate the ideal charge distribution in the

molecules to the observed binding energy, thereby giving an independent criterion for

alignment of the observed spectra. This is especially useful for examination of change in the

core level for the first adsorbed monolayers, where the core level shift can quantify the

charge transfer.

Based on this type o f description, the absolute binding energy is determined by the

amount o f charge on the atom o f interest, influenced by the number o f substituents, the

electronegativity o f those substituents and the formal oxidation state o f the element. All of

these factors must be taken into account in the prediction of the overall charge, q, for the

empirical equations to effectively predict the absolute binding energy. The best estimate of

this charge distribution is from ab initio methods based on Hartree-Fock density functional

calculations [6,7,8]. However, in the absence of such measurements, it is sometimes possible

to use empirical electronegativity equalization methods, such as the Jolly and Perry method

[6]; the Folkesson-Larsson method [5]; the Sanderson [9], or Modified Sanderson method

[10]; or the partial or full equalization of orbital electronegativity (P/F EOE) [11], Care must

be taken using these empirically derived equations, with the appropriate equation chosen for

the appropriate method of charge determination [7],

3.2.2 Extension o f Seigbahn theory

The attempts to establish the absolute binding energy for a given element as described above

generally neglect any final state effects. The effects of relaxation, however, often greatly



affect the experimentally observed energy of the core level and may be difficult to predict

theoretically. The absolute binding energies calculated from the various empirical or

theoretical formulae can often be unrealistic and impractical. However, it automatically

follows from the point charge model that the relative binding energy (the value o f the

chemical shift) due to a change in the local environment would neglect the zero point energy,

and may be expressed by

AEb (A ) = kjAqA + AVAMad + AER (3-5)

where AEb is the change of the core binding energy versus a reference compound, kj is the interaction coefficient between core electrons and valence electrons, A ja is the difference in the effective local charge on the atom of interest, AVMad is the difference in the Madelung potential due to the surrounding atoms if the sample is in the solid state, and AER is the difference in the relaxation energy term due to photoelectron emission.

In the classical theory [3], the Madelung potential can be approximately described by

the Coulomb interaction assuming each atom as a point charge in space. The point-charge

model for the potential in eV on a given atom, i, can be described by

F " = 1 4 . 4 a ^ — (3-6)i * j r ij

where otMad is the Madelung coefficient for a given crystal structure, qj is the charge on the other atoms, and ry their interatomic spacing in A 1.

The most practical application o f this type o f analysis still requires a good estimation

of the relative relaxation energies. The modified Auger parameter, described in section 3.3.2,

can be used as a basic estimate of this value. By combining such as estimate with an ab initio

estimation of the overall charge, the semi-empirical point charge model can give an adequate

approximation of expected binding energy shifts.

1 Note that the factor 14.4 in the equation o f the Madelung potential is to account for the numerical constants necessary to have the energy in eV.



Even with these simplifications, however, the interpretation of absolute chemical

shifts can be complicated by the possible coexistance of a number of factors not accounted

for by Seigbahn’s theory. Charging, analyzer work function effects, and alignment processes,

for example, do not readily allow for comparison of various systems. In contrast, the binding

energy splitting between two core levels can be used to examine the chemical state and

bonding configuration independent o f many o f these complicating factors since they

generally affect core levels o f different species equivalently [12]. This constitutes an

independent method o f examining the binding energy shifts to determine whether they are in

fact due to true changes in the chemical environment.

3.3 Use of secondary effects for analysis

3.3.1 Shake up features

In photoelectron spectroscopy, the photocurrent results from the excitation o f electrons by an

electromagnetic field from the initial state i, with a wave function \|/{ to the final state f with

wave function \|//. In this final state, the creation o f a core hole leaves the atom in an excited

state. If the photoemission process were slow, the electrons from or near the excited atom

would have sufficient time to change their energy by slowly adjusting to the effective atomic

potential in a self-consistent way [13] and relax down to the ground state o f the ion (adiabatic

approximation). However, as photoemission processes with X-ray excitation occur on the

scale of 10"17 s [14], there is no time for the ion to relax fully and there is a finite probability

that the ion is left in an excited state during photo-excitation with the creation of a core hole

[15,16]. In such a state, the wave functions o f the final state are a combination of the

ionization wave function and the wave functions o f the excited electrons, as in figure 3-4,



which compensate for the relaxation. The process o f excitation of electrons to a bound state

is referred to as a shake-up process, since it excites an electron into a higher orbital. If the

electron gains enough energy to be excited to the continuum, then the electrons are said to be

shaken-off [17],

T(i) ¥ (f)= 'P (io n )+ a 'P ( 1 )+ p T (2 )....

hv -

—o -

Fermi Level .........

-©— 4 —O

Intra-atomicexcitation

— © ---------------© —

Neutral atom — O-

Ek

Eb

Excited state ionFigure 3-4 Orbital redistribution due to the formation of a core hole

For most theoretical treatments, \|//, is considered to be decoupled from the initial

wave function such that it can be represented by the (N -l) electron system, with the loss of

both an electron and an orbital level. This is referred to as the sudden or frozen hole

approximation, [18]. If the relaxation is completely accounted for by electron excitation, then

the remaining (N -l) electrons maintain the same spatial distribution and energies in the final

state as they had in the initial state (frozen orbital approximation). In such a case, the binding

energy equals the negative orbital energy of the emitted electron, in Koopman’s theorem [2],

However, this requires excitation of all low lying electrons to their original orbital energies in

the neutral atom with the ejection of a photoelectron to accommodate the relaxation. Instead,

the core hole is usually screened by inter and intra-atomic electron energy redistributions and



the Koopman’s energy is never observed. The spectrum, therefore, is made up of an intense

peak from the original ejected photoelectron, shifted by the screened relaxation, and a

number of peaks on the high binding energy side o f the main peak due to the energy loss for

excitation o f electrons, as shown schematically in figure 3-5.

Main line

shake-up | Interatomic screening relaxation shifts

Adiabatic BE Koopman's Approximation Approximation{total relaxation) {no relaxation)

Figure 3-5 Schematic of the XPS spectrum under the sudden approximation

The peaks that have high binding energy from the loss due to excitation o f valence

electrons into a bound state on the atom are referred to as shake-up satellites. These satellites

can be thought of as reflective of the valence band characteristics, as the energies o f the

excited states of the (N -l) electron system can be approximated by ground-state energies

with valence electron promotion to an empty state.

The shake-up intensity in the spectrum is given by the projection of excited state,

with valence electron excitation on the frozen hole state. If the wave function o f the frozen

core hole state (with no valence electron excitation) and that of the relaxed ionic ground state

overlap, most of the intensity of photoemission will go into the main core level line [4], and

the satellites will not be well resolved. However, if the overlap is not high, then the intensity

of the excited ionic final states would be large enough to have them appear as satellites in the

photoemission spectra at higher binding energies [4], The shake-up excitation to unoccupied



states is favourable in a system if it helps to transfer charge to “screen” the core hole and

stabilize the ion [4], An intense shake-up satellite structure indicates a large charge

redistribution in the corresponding excited ionic state, or the presence o f convenient

unoccupied states within the valence band for electron promotion for optimum core hole

screening [19], as seen in conjugated organic molecules and in transition metals oxides with

partially filled J-orbitals. For organic molecules, the size o f the aromatic ring system has a

very significant influence on the satellites intensities [4], The satellite features have been

used widely for identifying changes in the conjugation when there is little obvious difference

in the C ls core levels for polymeric systems [20].

3.3.2 A uger excitation

Although the creation of a core hole causes reorganization and some relaxation o f the

electron energies in the surrounding orbitals, the resulting (N -l) system is still ionized from a

low lying orbital. In order to minimize energy, higher orbital electrons instantaneously relax

down into the created core hole [21], Often the energy gained through this relaxation is

sufficient to eject another electron, referred to as the Auger electron, as in figure 3-6. Thus,

photoionization normally leads to two emitted electrons - the core level and the Auger

transition.Auger electron

1-2,3 or 2p

Li or 2s— 0 - 0 - 0 - 0 - 0 0

0-0

10—0 K or 1s

Figure 3-6 Auger emission process [after 21]



As the photon energy changes, the kinetic energy o f the ejected photoelectron will depend

directly on the ionizing photons; however, as the Auger transition is related to the relaxation

o f electrons between orbitals, the kinetic energy o f the Auger electron is considered

independent of the initial excitation energy, and reflective of the relaxation occurring within

the system.

3.3.2.1 Relaxation and the Auger parameter

The definition of the total relaxation energy involved in the creation o f a core-hole state is the

difference between the energy from Koopman’s theory (frozen orbital approximation) and

that of the adiabatic approximation [22], as defined in the previous section. This relaxation

energy is a sum of different relaxation processes, including that of core electrons, valence

electron and any extra-atomic contribution to the relaxation, as in core hole screening. In

general, the relaxation energy of the core holes will be independent o f the chemical state,

while those of the valence electrons will be affected by any possible charge transfer from the

neighbouring ligands. In the Auger transition, if the Auger decay can be completely

separated from the creation of the initial core hole, the system is fully relaxed around the

primary hole [22], Similar to the definition for the photoionization peak, the relaxation

energy is the difference between this two step model and one where the Auger decay starts

from a non-relaxed initial core-hole state. As the binding energy for the core levels and for

the Auger peak shift equally with changes in the chemical state [22], the relaxation energy

can be defined by the summation of the energies of the core level and Auger level by

A(Ek (A u g er) + E b(core level)) = 2AE Reacore]eve] (3.7)

This summation of the core level binding energy and the Auger kinetic energy is referred to

as the modified Auger parameter:



a'=Ek(C' C"C"')+Eb(c) (3-8)

where C’,C” ,C” ’ are the Auger transitions for the core level excitation from the C orbital.

This parameter also has the advantage o f being independent o f surface charging (see section

3.4) and work function effects, making it useful for identifying chemical states for insulators

and for compounds with small core level binding energy shifts [23]. However, the Auger

parameter is only a one dimensional quantity, like the binding energy of the core level alone.

A more useful general approach is to display the photoelectron binding energy and the Auger

kinetic energy in the form of a scatter plot, called the Wagner plot, as shown in figure 3-7 for

Cu [22]. The position of the compound on the plot will be affected by both the initial (charge

distribution and local potential) and final (relaxation) state effects. On the Wagner plot, the

Auger parameters represent the intercepts o f lines with a slope o f one [22], All compounds

which lie along these lines have only initial state contributions in the chemical shift.

Compounds with similar initial state effects will appear on straight lines with a slope o f three

[22], following from the definition of the binding energy as in equation 3-5 above.

a'=1852.1 eV

CuCN

Figure 3-7 Wagner plot for copper showing the Cu 2p3/2 binding energy and Cu L}M4SM4S Auger kinetic energy for different chemical states. The straight lines with slope -1 represent compounds with the same Auger parameter, while those with slope -3 represent those with the same initial state effects. The binding energy and the Auger kinetic energy are referenced to the adventitious C Is line, set at 284.8eV[22].

937 936 935 934 933 932 931

Eb (2P*V eV



3.4 Charging in X-ray photoelectron spectroscopy

One of the side effects o f exciting electrons with irradiation is the build-up of a positive

charge density within the sample. If this charge density is predominantly made up of trapped

charges, they accumulate within a layer adjacent to the surface. The extent o f this surface

layer is related to the electric field that develops within the irradiated area due to the presence

o f holes. In the experimental conditions common to XPS measurements, the applied field is a

function of the trapped charge distribution in the material and the image charges in the

vacuum [24], For a thick planar film irradiated with a uniform X-ray flux, the solution to the

one-dimensional Poisson equation,

4 ? + — = 0 (3-9)dz £

where V is the potential, z is the distance to the analyzer parallel to the surface, Q+ is the positive charge density, and 8 is the material permittivity,

takes the form of a plane capacitor if the thickness o f the system is much greater than the

charged layer thickness [25], Therefore, the surface potential can be related to the developed

charge density by,

Vs ~ ^ h (3-10)£

where h is the thickness of the layer and Vs is the observed potential at the surface.

This build-up o f positive charges at the sample surface retards outgoing photoelectrons and

increases the apparent binding energy [26], as shown in figure 3-8(b), assuming a linear

relationship between the surface potential and the measured binding energy [27].



X-ravs

Photoelectrons

S P E C T R O M E T E R SAMPLE

Figure 3-8 (a) Schematic of charging inside a semiconductor or insulator during XPS measurement (b) Binding energy modification due to positive charging in the sample electronically decoupled from the spectrometer (adapted from [27])

As a result, equations 3-1 and 3-5 defining the binding energy of the outgoing electron must

be modified to accommodate this perceived change as

AEt (A) = k,A ,,A+ A V ^ - A E * +eV,(t) (3-11)

where e is the fundamental electron charge and Vs(t) is the surface potential as a function o f the irradiation time.

However, in good electrical conductors, the positive holes are quickly neutralized by the flow

of electrons from the surroundings, and very little change in the spectrum is observed. In

insulators, however, the lack of delocalized conduction band electrons within the sample and

the poor conductivity of electrons from the substrate can actively trap these charges leading

to a substantial charged layer adjacent to the surface [28], The manifestation of this effect in

the XPS spectrum, therefore, results from the competition between the emission of

photoelectrons into the vacuum, relaxation processes, and the electron redistribution from the

surroundings that imperfectly compensate for this emission over time [27,29,30], Even for

good conductors, however, the neutralization of charge is not instantaneous, and may



sometimes appear in the spectrum, especially if the sample is not properly grounded to the

spectrometer. For most samples, the charge distribution reaches a steady state within the time

frame o f the experiment, and any charging effects can be compensated for through alignment

with an internal standard. As the surface contamination layer shares the potential o f the

surface of the sample, adventitious CIs or OIs are good choices for calibration when

examining charging behaviour [31]. Some care has to be taken when analyzing the absolute

binding energy values and determining the chemical core level shifts after such alignment, as

the standard value o f C ls, 284.8eV, does not hold for some substrates [32], such as oxidized

or pre-treated metals.

Due to the need for charge movement, however, insulating materials will exhibit

some transient charging related to the compensating effects o f electron flow. At any time

during the irradiation, charge conservation has to be satisfied, and the movement of charge

within the material can be described by an equivalent circuit description [25], where

0 R

f \? (3-12)

v dt/

where I0 is the secondary electron emission current, Vs is the surface potential and R is a resistance value incorporating both sample resistivity and the self-regulation effects from the vacuum as described in section 3.4.1 below, S is the uniformly irradiated specimen surface

area, and is the algebraic rate of change of the charge density.

The secondary electron emission current is related to the secondary electron yield, S , by

I c = e9 S S '(V s ) (3-13)

where 0 is the X-ray irradiation flux.

Substituting equation 3-13 into equation 3-12, the time dependence of the surface potential,

and therefore, the binding energy, can be described by



d r s t Vs F S e 5 ‘ (Vs )dt RC C

where C is the geometric “capacitance” defined by the ratio between Vs and Q+S.

If the secondary electron emission is considered independent o f the surface voltage, the

solution to equation 3-14 would take the form of an exponential decay [24, 27],

- / / \

where AFX00) is the steady state value of the binding energy shift, and r is the effective time constant, x=RC.

Generally, however, the secondary yield is not independent o f the surface potential, and the

determination of R and C relies on the initial state and final boundary conditions for equation

3-15, as

dt Cat t=0 , and

3 r , _ F SqF(o\ (3_16)

R

at t= 8 .

F S q S x{Vs ) (3-17)

3.4.1 Charge compensation

As described above, the charging response of the XPS spectrum is very complicated. There

are many possible charge compensation mechanisms beyond the conduction o f electrons

from the sample and its surroundings through the sample itself [29,33], The magnitude of the

charging is generally more a function of the experimental conditions and of the crystalline

state of the specimen than its exact chemical composition [29]. Table 3-1 lists the many

parameters that can affect the charging behaviour o f a given sample.



Table 3-1 Charging mechanism parameters (adapted from [29])Factor Contribute to charging through Compensation for charging through

Composition:

Relative dielectric constant of the material, er; the secondary electron yield of the specimen, 8X; Photoelectron cross section

DC conductivity of the specimen

Crystalline state: 8X and the trapped charge density, ptrTemperature P*Surface composition /cleanliness

Effective retention coefficient or electron affinity of surface, 8X ptr

ThicknessPhoto radiation induced conductivity (RIC), photoelectrons injected from substrate

Specimen /holder contact

DC conductivity; RIC; lateral compensation from clips, metal grid

X-ray tube window Flux of electrons from surroundings

Grounded parts Secondary electrons induced by photoelectrons

Flood gun Flux of electrons from surroundingsVacuum pressure and composition Flux of electrons from surroundings

Ion gauges and pumps Flux of electrons from surroundingsIncident flux 8X Bremsstrahlung radiationIrradiated areas 8X Flux of electrons from surroundingsPhoton energy 8X Photo radiation induced conductivity

Sources of charge compensation are also shown schematically in figure 3-9. If

experiments are performed in the same chamber under similar conditions, charge

compensation from the surroundings and the irradiation conditions is fairly consistent.

Compensating electrons from the surroundings would be virtually unchanged for all

experiments where the same acquisition gun and sample holder are used in the same chamber

at approximately the same pressure; therefore, any electron or secondary electron emission

due to [29] the gun material, the chamber walls, the X-ray window (not applicable for

monochromator use), the gauges and pumps, the vacuum pressure and composition would be

the same for many cases. Other factors [33] that may affect the charging behaviour are the



photoelectron cross section, the effect o f electron emission from the solid on the effective

attenuation length, the electron yield (electron emission distribution per photon), and the

effective retention coefficient/electron affinity of surface.

Figure 3-9 Sources of charge compensation. Left: sources issued from the specimen holder. Right: sources issued from the surroundings of the specimen that are normally to ground. The dotted lines indicate the incoming X-rays, and the solid lines the electron movement in the sample (adapted from [29]).

3.4.2 Use o f charging fo r electrical information with XPS

With the knowledge o f the charge compensating mechanisms possible in XPS, it is possible

to use the sensitivity to the conductivity o f the films under irradiation to gain some

information about the relative conductivity of thin films. Traditional conductance techniques

may not be suitable to accurately measure the conductivity of thin organic films and

molecular devices [34,35]. The introduction of any “external” contact to the film of interest

to probe the properties will effectively change those properties such that isolating the

electrical behaviour o f the film becomes very difficult. XPS, on the other hand, can be used

as a non-contact method for analyzing the resistance/capacitance and other electronic

properties of thin semiconductor and dielectric films [27,35,36,37] by examining the effects

of charging on the position o f the core levels (surface charge spectroscopy [37] or chemically

resolved electrical measurements [35,38]). An estimate of the relative electrical properties of

insulating films on different substrates irradiated under similar conditions is possible by

examining charging effects using basic electrostatic theory from equations 3-10, and

Secondary electrons from w alls or detector

lateral 'compensation (i.e. h o i'd e r y ')-* — Radiation induced

f conducth'hv Electrons/ions from vacuum

Electrons from flood gun

Electro nssfrom window *

Photoelectrons injected from substrate

Electrons from substrate/ground

lateral com pensation

bad contact



equations 3-15 to 3-17. Refinement of the estimation is also possible by changing the X-ray

irradiation conditions [27], though this was not done in this thesis. In any case, care must be

taken when interpreting the results as the parameters o f such a technique are not fully

established, and the values derived are not necessarily reflective o f the absolute conductivity.

3.5 Angle resolved XPS

3.5.1 Information depth o f photoelectrons

Though all photoelectrons are emitted with an angular distribution, the angle o f detection

determines the depth from which the emitted photoelectrons may be observed. From the

diagram in figure 3-10, it can be seen that detection close to the surface normal enhances the

signal from the bulk relative to the surface, while detection close to the surface plane

enhances the signal from the surface relative to the bulk.

Detector

Detector

Figure 3-10 Schematic of angle resolution

If X is the effective attenuation length (EAL) o f the emerging electrons, then, for take-off

angles normal to the surface, 95 % of the signal intensity is derived from a distance 3X in the

solid, assuming an exponential energy loss through the material thickness as in figure 3-11.



0.69-

©

0.2

M£0 20 40 m 80

Depth of creation d, A

Figure 3-11 Exponential decay of lossless electron escape with depth of creation of photoelectron [39]

The vertical depth sampled is given then by:

d = 3A, sin 0 (3-18)

where A is the effective attenuation length, and 6 is the electron take-off angle, with respect to the surface plane

Thus, varying the angle o f the sample surface relative to the detector can change the

measured excitation region.

If the films are not homogenous through the thickness, the composition at depth can

be determined nondestructively, as shown for SiCE on Si in figure 3-12. This approach is

obviously preferable to destructive ion sputtering, but its applicability is limited by a number

o f assumptions regarding the physics o f angle resolved XPS (ARXPS) measurements [40].

Primarily, the specimen must be amorphous or finely polycrystalline with an atomically

smooth surface, the electrons should undergo minimal elastic scattering with no refraction at

the sample surface so that the exponential decay of signal is observed, the EAL should be

independent of composition, the acceptance angle o f the detector should be small, and the

algorithm used should cope with a wide variety of poorly resolved peaks o f varying intensity

without introducing systematic errors. Most of these conditions are easily met for simple

analyses o f compositional changes.



Si 2p

Grazing

Normal

Figure 3-12 Enhancement of surface composition in core level intensity at grazing angles for Si 2p [211.

UO 96B inding Energy (eV)

3.5.2 Thickness and coverage dependence o f overlayers

In addition to compositional analysis, the well-known relationships regarding the attenuation

of photoelectrons with changing angle o f acquisition can be used to reliably estimate the

structure and thickness o f overlayers [41]. In figure 3-13, the expected response to varying

the take-off angle for a variety of overlayer configurations is shown, along with the defining

equations.

As can be seen from figure 3-13, for thin overlayers, thickness can be estimated from

the ratio o f the intensity of a core level from the overlayer and one from the underlayer, if the

appropriate EALs are known. For simple surfaces, where the overlayer and underlayer share

a common element, assuming a uniform film forms on the surface (uniform overlayer model

[41]), the thickness can be estimated from [42]

'k t o underlauer J -/V 1 overlayerd = XA sin 9 In

v r o overlayer t ^ B ^ B 1 underlayer

+ 1 (3-19)

where Iunderiayer and Ioveriayer are the intensities of the photoelectron peaks from the underlayer and the overlayer respectively, Nb and Na are the densities of atoms in the overlayer and underlayer (atoms/cm3), y f erk<ycr and y ^ rIawr are the EALs of the photoelectron through the underlayer and the overlayer material respectively, and 9 is the electron take-off angle, defined by the surface plane.



A

} < 5 0 A

A

) < 50A

} > 50 A

(// ............

\ \ v

ta k e -o ff a n g le (6 )

take-off an

' //

gle (0)

take-off arig le (0) -

7 7 = exP1A Aa sin 9

h ,-f- = 1 - exp I R y 2s sin Oj

1 X exPd

A , sin 91 - exp -

X„ s in #

ta k e -o ff a n g le (0 )

Figure 3-13 Response of various overlayer configurations to changes in the take-off angle for photoelectrons.

For overlayers where a coherent film has not formed, the equations in figure 3-13

may be used to estimate the extent of overlayer coverage on the surface. If the overlayer and

underlayer do not share a common element, an island overlayer model [41] can be used

taking a ratio o f the intensity of the overlayer peak to the underlayer peak.

I°A

1b

n

= X expv XA sin 6

■■x 1 -e x p

(* - * )

/II sin 0

(3-20)

J)where d is the thickness of the overlayer islands, and x is the area fraction of coverage.

Modifications of these simple overlayer models are possible, but care must be taken

in interpreting the results. ARXPS has the smallest degrees o f freedom of any existing

remote sensing technique [40], Any model assessing the thickness or coverage can have no



more than three adjustable parameters to give any meaningful solution from the

measurements alone, regardless of the number o f angles chosen for measurement. Even with

this few adjustable parameters, the precision of the peak intensity measurements must be less

than 3% for the retrieved values to determine a realistic picture of the film structure [40]. A

fit of simple parametric models containing only one or two adjustable parameters per element

is likely to give more accurate values of those parameters than if they were measured from a

depth profile plot obtained from any general inversion algorithm. Due to these constraints,

ARXPS when applied to structure or thickness calculations generally cannot also be used to

determine the composition of the overlayers as well as thickness and coverage.

3.6 Equilibrium chemical states analysis of interfaces in OLEDs with XPS

As mentioned in chapter 2, both the performance and the reliability o f devices are intimately

linked to the chemical and electronic state o f the interface. The performance of devices with

a variety o f different metals, alloys and interlayer systems indicates that generally much

remains unclear about the role of the chemical state for different cathodes in controlling

device performance. Specifically, the superiority of Mg:Ag and LiF/Al cathodes and any

possible link to the chemical structure of the interface is yet to be fully explained. Many

investigations were undertaken to examine the states formed due to the interaction of the

organic and the cathode. Since the region of interest is limited to the interface between the

two materials, a variety o f surface analyses have been employed, XPS being the most widely

used technique. A majority of groups analysed the interface formation process by growing

metallic or organic overlayers on a suitable substrate, until the signal from the substrate had

been obscured. Only a few researchers examined the buried interfaces directly [43,44].



The literature indicates that there are predominately two interface types, based on the

cathode used. Non-reactive cathodes, such as Ag [45,46], Au [46], Ga [47], show the

formation of a diffuse interface, with the cathode metal diffusing well into the organic layer

without any reaction. For reactive cathodes, there is a general consensus that some type o f

chemical interaction is occurring between specifically Alq3 and a number o f cathodes,

leading to the formation of an interfacial reaction zone. However, the interpretation of that

interaction upon deposition is controversial, especially with regards to Mg and LiF/Al, the

most widely used cathodes. Some groups claim the interaction is dominated by the formation

of a radical anion with either a primary charge transfer interaction with N of the pyridinal

ring or with the O of the phenoxide ring [48,49,50,51,52]. Other groups suggest that there is

a fragmentation of the molecule with the formation o f an oxide or an organometallic

compound [46,52,53,54,55,56]. There is some evidence supporting all o f these possibilities,

as described below, preventing a final conclusion about the dominant interaction mechanism.

3.6.1 Low work function metal cathodes

A number of studies have been conducted into the behaviour of low work function metals at

the Alq3 surface, since the earliest accepted models of injection suggested that a low work

function was a requirement for acceptable performance. The first systematic photoemission

study of such an interface was with the alkaline earth metal Ca [53,57], XPS analysis showed

that with deposition of more than 4A, the O Is core level, which was a one component

Gaussian for pure Alq3, split into two peaks, as shown in figure 3-14 below. At the same

time, the Ca Is core level showed a similar chemical shift. Therefore, Ca was thought to react

destructively with Alq3, forming an oxide with deposition o f more than 4A of Ca, visible as

the new chemical state on the high binding energy peak on the O Is.



O 1s

528 526532 530534

Binding Energy (eV)

Figure 3-14 Deconvoluted O Is core level [53]

At lower coverages o f Ca, where the Ca seems insufficient to initiate a reaction with O,

Choong et al. [53,55] observed the emergence o f a low binding energy shoulder on the N Is

core level suggesting that the deposition of metal causes a transfer o f charge from the metal

to the N in Alq3, before the onset o f molecule fragmentation. With the emergence of the new

oxide species, the split peak in N remained at constant intensity as the thickness increased,

while the O species continued to grow. This evolution of peaks indicates that the Ca transfers

charge to the N peak until saturation occurs, with two ligands accepting charge, forming a

stable anion.

Density functional theory calculations [58] on charge transfer to Alq3 indicates that a

transferred electron would be localized to the pyridinal side o f the ring, presumably at the N

linkage, which has the weakest bond in the molecule. The valence spectra, observed by UPS,

showed the emergence of states within the HOMO-LUMO gap even with low coverages of

Ca presumed to be from the Ca-N interaction, and a shift in the orbital structure of Alq3 to

higher binding energies as would be expected for atom interaction with extended % electron

systems [48]. These were taken as further evidence of the formation of a radical anion. Work

2+by Curioni et al. [59], however, indicates that such a charge transfer, with Ca stabilizing the



radical anion of Alq3, is not thermodynamically feasible. The most likely site o f interaction

to form the radical anion is actually on the O on the phenoxide side o f the ring, prior to

molecular fragmentation. Photoemission spectroscopy o f the reverse deposition with Alq3 on

Ca immediately showed the emergence o f oxidized states and composition ratios inconsistent

with Alq3, indicating that the destructive diffusion limited reaction dominates the interaction

o f Ca and Alq3.

Further XPS work with Li[60,61], K[62, 63], and Na[54] also appeared to support the

formation o f a radical Alq3 anion upon deposition, but without the destructive oxidation

reaction. All three metals showed the emergence o f a lower binding energy state in the N Is

core level and the appearance of gap states and a shift in the binding energies in the valence

band. However, in those cases, the new N species continued to grow with deposition. For Li

and K, the peak grew until the ratio o f metal atoms to Alq3 molecules was 3:1 where it was

again a single component Gaussian peak, shown in Figure 3-15 for K, indicating that all N in

the molecule have gained charge.

. . . . A \ .

K/Alq3 * 3

c3

Figure 3-15 N Is evolution with K deposition [60]

PRISTINE Alq3

410 408 406 404 402 400 398 396 394

BINDING ENERGY (eV>



In contrast, Osada et al. [61] observed the formation of the new species in the N Is

core level on the high binding energy side for K deposition. They found that the difference

between the energy of the new species with Li deposition and that with K deposition was

equal to the difference in the work function o f the two metals. Although this seems to suggest

that K was drawing charge away from the molecule, Osada et al. believed that both Li and K

were beneficial dopants in the organic. Theoretical calculations [59] on Li interaction

suggests that in this case the N bond is feasible, but that the O should again be the primary

interaction site. Osada et al. [61] claim that there is no effect on the O Is core level with K

and Li deposition, though the theoretical calculations on bond length by Johansson [60] and

Curioni [59] indicate that there should be some interaction on the O on the phenoxide ring.

Although Li and K are not widely used as cathode materials, Li doping o f the organic

layer through co-deposition has been shown to improve the device performance, indicating

that there can be some beneficial charge transfer between metal and organic creating new

carriers, presumed to be the radical anions [62].

3.6 .1.1 Mg cathodes

For Mg, the experimental results are even more varied. A number o f groups, using UPS, [49,

50,51,52] reported the existence of a gap state in the valence structure, as well as a shift in

the HOMO and vacuum level [49,51,63,64], This behaviour is similar to that observed for

other metals as shown in figure 3-16, indicative o f the radical anion formation discussed

above. He et al. [50] also saw evidence o f reduced symmetry o f the molecule due to Mg

attachment, through the broadening and shifting of the features in a high resolution electron

energy loss spectroscopy (HREELS) study upon deposition of Mg on Alq3.



1.6 eV

H*z=s Figure 3-16 UPS spectrum for the

HOMO region of Alq3 after deposition of various cathodes [52].

s<>-t»zLU

Na

I-Z

Ca

21 0 S 6 4Binding energy (eV)

Theoretical investigations o f Mg and Alq3 have indicated that the likely site o f

interaction is again at the O in the phenoxide ring [65], though Curioni et al. [59] claim that

the low affinity of Mg for Alq3 would limit any radical anion formation. They claim that the

use of semi-empirical techniques by Zhang gave a poor representation o f the possible

structures, leading to a misinterpretation of the more energetically favourable orientation.

Although the bonding is thought to be quite weak, it was presumed to be sufficient to

develop the gap state observed in the valence structure.

Both the Princeton group [54] and the Kodak group [52] reported the emergence o f a

low binding energy shoulder on the N Is core level, also indicative o f radical anion

formation, with Mg deposition on Alq3, though showing much less intensity than that

observed with the monovalent species. Mason et al. [52] attributed this difference, also

observed with Ca, to the ability of each species to give up charge to the molecule. They

suggested that given a sufficient supply o f metal atoms, alkali metals are able to donate up to

three electrons to each molecule, whereas the alkaline earths can only donate one electron to

each of two Alq3 molecules. However, the interpretation of Shen [54] that the reaction with

Mg is more likely to form an organometallic product characterized by Mg-C bonding, and



our results [46] showing oxidized Mg and metallic Al at the buried interface in real devices2,

as well as the results o f Choong [55] on Ca, could indicate instead that the destructive

reaction of Alq3 supersedes any charge transfer. Since reaction products could have valence

structures that differ from the reactants, the gap states observed by UPS can also be attributed

to the reaction products themselves.

Investigations on the degradation o f devices with Mg based cathodes show some dead

spots that correspond to fully oxidized cathode materials. Examination o f the buried interface

indicated that severe chemical reactions had taken place [43,44]. It is worth noting that these

dark spots occur upon deposition, and no new dark spots are formed during the operational

lifetime of the devices. As well, He et al. [50] reported the formation of Mg-O stretch mode

in the HREELS spectrum only after heating the substrate. Therefore, the Mg may require a

catalyst or trigger, such as oxidation o f Mg with external oxygen as could be encountered

during conventional fabrication, to initiate further reaction with Alq3. Since there is little

experimental information in the literature about the activation energy required for reaction,

there are still many unanswered questions about the interaction of Mg and Alq3.

3.6.2 Bilayer cathodes

3.6.2.1 Al and Al/LiF cathodes

Al/Alq3 and Al/LiF/Alq3 interactions have also been the subject o f both experimental and

theoretical investigations. With Al deposition, again a low binding energy shoulder was

observed on the N Is core level [52,56,66], though some groups reported only peak

broadening [57,67]. Although this is one o f the indications of radical anion formation, Al is

considered to have a destructive reaction with Alq3, shown by the broadening o f the Al 2p

2 Explained in greater detail in chapter 5.



core level [52], the formation of a high binding energy shoulder on O Is [52,56,66], and the

elimination of all features from Alq3 in the valence band after as little as 0 .2-1 A deposition o f

Al [51,52,56,57,67]. Theoretical calculations indicate that the most likely site o f interaction

is at the O, as in Li, K, and Ca [52,59,65],

LiF deposition, on Alq3 or on Al, is seen to have no reaction, with all the core levels

showing single component Gaussian peaks corresponding to Alq3, Al, or to LiF [56,68,69,

70,71], and no effect on the valence structure. This has also been observed for MgF2 [70,72]

and CsF [52], though CaF2 did show a HOMO shift in the UPS spectrum [70]. Upon

deposition of Al on top o f the fluoride, the observed spectrum is different than both that o f Al

by itself and of the fluoride by itself, showing the emergence of a gap state below the Alq3

HOMO [52,69,70,72]. For all fluorides, subsequent Al deposition also produced a N f r low

binding energy shoulder [52,56,70,72]. Since these two features were also observed for

radical anion formation with low work function metals, it was theorized that deposition of Al

caused the LiF to dissociate, thereby doping the organic and producing a radical anion [52,

56,68,69]. As this dissociation mechanism could also be used to explain the improved

performance with a LiF interlayer, it has become the most popular interpretation o f the role

o f LiF in devices. Mason et al. [49] suggested that one possible dissociation pathway was

through the formation of AIF3, which, though thermodynamically unfavoured directly, may

be forming due to the presence of all three species, Alq3, Al and LiF, together.

However, none o f these studies, nor investigations with other techniques [73], ever

observed Li or F in any state other than that of LiF. In these investigations, the deposition of

metal atop the LiF completely obscured the signal from Li, since Li has a very low

photoionization cross section. However, there was also no report of any changes for the

observed F Is core level, which has a much higher photoionization cross section. The



HREELS investigation o f Hung et al. [73] on Al deposited on LiF/Alq3 indicated LiF

dissociation indirectly through the attenuation of the Li-F stretch mode with Al deposition,

but again did not report the emergence o f another Li stretch mode showing the subsequent

behaviour of Li in the system. Subsequent investigations of the LiF/Al interaction have

shown no evidence o f AIF3, either by deposition onto Al [74], or with other organic

molecules [75], Grozea et al. [76], indicate that the F Is core level has a high binding energy

shoulder.3 Although this could be an indication of LiF dissociation, the Li Is core level for

the systems that would have confirmed this phenomenon was not observed. The F Is core

level with C6o/LiF interactions, described in section 8 .2 , also have a shoulder at the same

position as that observed in the Alq3/LiF system, suggesting that the F shoulder is due to a C-

F interaction, rather than AIF3 formation. Since F' ions would likely react with the metal

cathode before the organic molecule, this raises a number o f questions about the supposed

doping of the organic layer with free Li.

3.6.3 Limitations o f previous studies utilizing XPS

Extensive work has been undertaken regarding the interface formation process, and though

there is some understanding of the interfacial interactions, there are still a number of

unanswered questions.

The investigations described above generally follow a classic surface science method,

with deposition of thin metallic or organic overlayers to examine interface formation. This

type of analysis may not give the complete picture regarding the structures that can be

formed in real devices, produced under conventional fabrication techniques.

3 Explained in greater detail in section 8.1.



Since conventional fabrication techniques rarely involve layer by layer deposition o f

the cathode over large time frames, this type o f analysis assumes that the steady-state

condition reached at every stage of the deposition during analysis is the same as that in a

conventionally fabricated device, discounting the quasi-static condition reached at each step.

The merit o f monolayer deposition would be to ascertain the reaction pathway and develop

an understanding o f the kinetics of interface formation, and the activation energy for

interface formation, which are not covered extensively in the literature. Secondarily, this type

of analysis is performed from the top down, limiting the deposited thickness to the probing

depth of XPS before the signal from the interface itself is obscured; as such, it is of limited

use in examining buried interfaces. Thirdly, as there is no way to separate the cathode from

the organic in such an analysis, it is difficult to draw conclusions regarding the chemical state

o f elements that the two materials have in common or the extent o f band bending at the

interface.

Some researchers have done sputter analysis through the thickness o f a deposited

cathode to expose the buried interface, either by XPS or SIMS, but the damage that occurs

with sputtering in Alq3 [77] can induce artefacts that can obscure the effects of interest.

Ca cathodes provide a good example o f the limitations o f the traditional surface

science techniques to gain an understanding o f the equilibrium chemistry at the interface of

real devices. With the onset of the destructive reaction between Alq3 and Ca occurring at 4A,

real devices will always have enough cathode atoms to attack the organic layer. Other

systems may show similar diffusion limited reactions at greater thickness where the

traditional method may no longer be able to probe the interface. There may also be other

kinetically limited diffusion and reaction processes occurring during fabrication of real

devices, which sometimes cannot be captured by model interfaces formed in the lab. This re



emphasizes the importance o f the analysis o f real equilibrium device structures in

conjunction with the standard analysis o f thin deposited layers. The novel application o f a

peel-off procedure, as described in the next chapter, to XPS extends the capability o f the

technique to the examination of the interfaces occurring in buried junctions, such as those

found in OLEDs.

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67 T. P. Nguyen, J. Ip, P. Jolinat, and P. Destruel, Appl. Surf. Sci. 172, 75 (2001).

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69 T. Mori, H. Fujikawa, S. Tokito, and Y. Taga, Appl. Phys. Lett. 73, 2763 (1998).

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Chapter 4

Experimental

4.1 Molecular beam deposition/Vapour phase deposition theory

Physical vapour deposition (PVD) is the growth of thin films through the production and

deposition of a condensible vapour by physical means [1]. The atoms or clusters of atoms in

this vapour are often those not normally found in the gas phase [2]. The objective of PVD is

to controllably transfer the vapour atoms through a chamber under high vacuum to a

substrate located a distance away. As the molecular beam impinges on the surface, film

formation and growth proceed atomistically [3]. The most widely used method for forming

such a vapour is by thermal heating of the source material, called thermal evaporation. When

the source is sufficiently hot, atoms either evaporate or sublime from the source and

condense on the substrate.

- 6 2 -


Chapter 4 Experimental 63

The classical “hard sphere model” [4] can often be used to describe the behaviour o f

thermally evaporated vapours, assuming that the molecular beam consists of an ideal gas

made up of infinitely hard spheres, undergoing only purely elastic collisions. With pure

element evaporation, the individual atoms themselves are modeled as hard spheres;

compound evaporation, however, can consist of different molecular configurations,

potentially with dissociation and molecular fragmentation [3]. For Alq3, the molecular

configuration o f the vapour phase is somewhat dependent on the evaporation temperature,

with some incorporation of nitrogen species from the vacuum modifying the observed

stochiometry at deposition rates below 1 A/s [5], Fullerene molecules, on the other hand, are

extremely robust, maintaining their cage-like structure during free evaporation [6 ] even at

extremely low deposition rates [7]. Many ionic solids, especially metal halides, tend to form

polymer molecules during thermal evaporation, rather than dissociating. LiF, unlike most

alkali halides, tends to form a small proportion of trimers, in addition to monomers and

dimers in the vapour phase [8,9,10]. For such molecular solids, the intact molecules may

themselves be treated as classical particles in an ideal gas.

Regardless of the nature of the vapour species, the evaporation rate from a heated

source can be defined as

J .=a . , P"r (4-1)V2mnkT

where Je is the evaporation flux in number o f atoms (or molecules) per unit area per unit time, (Xe is the evaporation coefficient, Pvap is the equilibrium vapour pressure of the source material at temperature T.

Thermal evaporation can be carried out in quasiequilibrium sources, called

effusion/Knudsen cells [1], or in nonequilibrium open sources, such as evaporation boats or



crucibles. All such sources can all be considered as flat, because evaporation is

unidirectional. In a quasiequilibrium cell, the evaporant is almost in equilibrium with its

vapour, and the particles escape through the opening at the mouth o f the cell in a collision-

free fashion. The cell is then said to produce a “molecular beam” or an effusive stream with a

Maxwell-Boltzmann velocity distribution in a molecular flow regime. For non-equilibrium

sources, there is no return of the evaporated vapour flux to the source [2 ], and all the material

undergoes free evaporation. Non-equilibrium sources, therefore, tend to have wide emission

angles, predicted by the cosine law of emission [1 1 ],

Effusion sources, on the other hand, are often specifically designed to deviate

somewhat from the ideal. Many commercial effusion sources have an open-tube

configuration that strongly focuses the molecular beam into small emission angles around the

axis normal of the cell. The improved directionality allows the use o f masks on or just above

the surface to delineate the deposition area. The drawback, however, is that for large

substrates, the increased fraction of arriving molecules with beam focussing reduces the film

uniformity. These effects may be accommodated somewhat by rotating the sample during

deposition [3].

As the emission beam reaches the substrate, the conditions at the surface control the

deposition process. The temperature and composition of the substrate have little impact on

the particle arrival process, but can strongly affect the mobility and residence time o f the

particles on the surface, which can strongly impact the resultant film properties [2 ],

Condensation on the substrate surface, therefore, is a function of the emission process minus

the effect of adatom revaporization. The flux o f atoms that condense can be described by a

variation on the Hertz-Knudsen equation [1], as



J q cos(p P Z (T sub)

( 4 - 2 )

where Oc is the condensation coefficient (the fraction o f incident particles that actually condense on the surface), Jq is the molecular flux as a function o f the angular distribution from the effusion source, q> is the emission angle, r is distance to the substrate, (% is the revaporization coefficient, is the equilibrium vapour pressure o f the source material above a substrate at temperature Tsub

For practical deposition rates, the temperature of the substrate should be kept low

enough that supersaturation of the molecular beam above the surface is possible and

condensation can occur. For most film growth with the substrates held at room temperature,

revaporization can be neglected and film growth is mainly controlled by the geometry o f the

evaporation system. Though the condensation coefficient is related to the energy and diffusion

processes occurring across the substrate, for most systems, Oc can be taken as unity [1].

True adsorption o f adatoms on the surface cannot occur without energetically

favourable positions for the vapour phase molecules. In order to become a part o f the

growing film, the impinging atoms must first be adsorbed, then migrate across the surface

and bond to another atom. Depending on the nucleation probability and surface energy, the

absorbed adatoms may desorb back into the beam or be reflected specularly off the surface.

The difference between purely van der Waals or electrostatic adsorption (physisorption) and

true chemical bonding between the adatoms and the substrate is related to the magnitude of

the desorption energy. The desorption energy, Ed, is defined as the depth o f the potential well

with respect to the energy of a particle infinitely far away from the surface [1], If Ed,

generally determined by the minima o f the integrated Lennard-Jones potential, is less than

0.4eV, the adsorption is classified as physisorption; if it is above leV, the particle is said to

be chemisorbed with strong chemical bonding on the surface.



Regardless o f the adsorption mechanism, the equilibrium adsorption at a specific

temperature can be described by Langmuir’s adsorption isotherm equating the rate of

adsorption to the rate o f desorption [1 2 ],

= N sX vex ) (4-3)

where Ns is the surface density of adsorption sites, x is the fractional surface coverage, v is the frequency o f vibration normal to the surface (usually taken as 1013 Hz), Ed is the desorption energy, and Br is the trapping probability.

The deposition process and properties o f the deposited film, however, is also affected

by the thermal power delivered to a substrate during deposition. Generally, the atoms or

molecules arriving at the substrate surface have a kinetic energy consistent with the

temperature of the deposition source [2 ], governed by the kinetic theory o f gases, yielding a

power density of:

^ k = J e^ k T (4-4)

where Je is the evaporation flux, and T is the source temperature.

This impingement energy has little impact on the deposition process since the arrival

energy for thermal evaporation is typically 0.05eV, with 0.1-0.15eV for evaporation above

500°C [2]. However, the condensation energy for the arriving particles can be substantial,

based on the sublimation energy of the source material:

W 'c = — AH c (4-5)N a

where AHC is the heat o f sublimation/condensation, and T is the source temperature.

For molecular films, this amount of energy is often sufficient to activate reactions between

the arriving molecules and the molecular film [13], leading to the “hot atom” impingement

driven chemisorption [14], This can also be aggravated by radiation heating from the



evaporation source, if the throw distance between the source and the substrate is short, as in

many traditional surface science experiments. The radiation power density at the substrate is:

= e s ( 7 s T s ( 4 - 6 )

r 2where es is the source emissivity, crs is the Stefan-Boltzmann constant (5.67x1 (T8 w/ m 1k4 ), 8Ais the area of an emissive source, Ts is the source temperature and r is the distance between the source and the substrate.

In this thesis, a throw distance o f ~20cm, as in commercial OLED fabrication tools, was used

to minimize the effects o f radiant heating at the substrate.

4.2 Instruments

4.2.1 MA C in-situ system

The major facility used in the course o f this project is the Multi-Access Chamber System

(MAC), shown in figure 4-1. The MAC system, assembled in the summer of 2001, consists

o f 5 chambers used for deposition (OMAC and NMAC), sample preparation (delamination)

and analysis (<f>-MAC). These chambers are connected together through the central

distribution chamber o f the MAC system (CMAC), which consists o f a stainless steel UHV

chamber with multiple ports to accommodate all the attachments. The base pressure of the

CMAC is kept around 10"9 Torr to ensure minimal sample contamination. Each attached

chamber is isolated from the CMAC by a high vacuum valve, allowing for reconfiguration as

necessary. Samples can be loaded into the MAC system through the analysis chamber (<D-

MAC), or through the load-lock situated on the CMAC. The load-lock is also used to store

ex-situ deposited samples at high vacuum (pressure ~ lx l O’6 Torr). The sample mounting fork

in the CMAC can rotate 360° and extend into the centre of any attached chambers to allow

movement o f the sample from chamber to chamber without breaking vacuum.



Figure 4-1 Multi-Access Chamber (MAC) System

Deposition o f organic or inorganic materials can be performed in one of the two

deposition chambers, referred to as the OMAC and NMAC for organic device and inorganic

device fabrication respectively. In this thesis, the OMAC was used for a majority o f the in-

situ experimental work, and is described in greater detail below. The NMAC is equipped

with a high temperature Knudsen type molecular beam source designated for inorganic

semiconductor materials.

4.2.1.1 OMAC

A majority of the in-situ experimental work was performed in the OMAC chamber. An 8 F

Cryotorr cryopump maintains the OMAC at a base pressure of 3xl0 ' 9 Torr, monitored by two

high vacuum glass encased external ion gauges from Kurt J. Lesker and Varian. The

substrate sample holder has an x-y-z stage manipulator with 360° rotation capability, and is

equipped with a home built halogen heating system for substrate outgassing prior to



deposition. The OMAC consists o f two evaporation source panels: a commercial deposition

panel for organic deposition, and a home built source panel for LiF deposition. Both are

situated at a throw distance of 2 0 cm from the substrate sample stage to minimize the effects

of radiant heating o f the substrate by the source. Due to the small size of substrates used in

this study, the stage was not rotated during deposition. Deposition amounts were monitored

with an Inficon XTM/2 oscillating quartz crystal microbalance, held parallel to the substrate

holder. Thicknesses were calibrated by both X-ray photoelectron spectroscopy and by mass.

A schematic of the OMAC chamber with ports and distances is available in Appendix B.

4.2.1.1.1 Cathode source

For inorganic material deposition in the OMAC chamber, a source flange was designed and

built. To maintain a reasonable throw distance to the substrate, four support rods 6 %” long

were welded to an 8 ” stainless steel flange. A 5” diameter steel ring was bolted to the ends of

the support rods, to form the basis for the electrical supports for the sources. The support ring

was used to allow maximum flexibility in source configuration, with the sources positioned

according to current needs for deposition. For most experimental work with the source

flange, a two source configuration was used with two support rods bolted to the sources and

to the support ring. The boat support was electrically isolated from the support ring with

ceramic spacers. Evaporation of the inorganic materials was performed by resistive heating

either a thermal boat source or a pyrolytic BN or AI2O3 crucible in a wire basket, as shown in

figure 4-2 below. The source was attached to the electrical support rods using Cu spacers.



C R U C IB L E W IT H B A S K E T

Figure 4-2 Resistive sources used for thermal evaporation of inorganic materials in the OMAC chamber [1]

To prevent thermal cross-talk between the two sources, Mo shielding was attached to

the non electrical support columns welded support ring. The final design is shown

schematically in figure 4-3. The temperature reached is estimated through the applied current

and the deposition rate, rather than through the direct use of a thermocouple.

Mo shielding

Cu contacts to crucible \

Cu contacts"

Crucible heater with basket10 crrj

Mechanical support

3.4 cm!

electricalinsulation

Electrical contact Mo shielding

! 12 .7 cim

] electrical control [

(a)

Figure 4-3 Schematic of cathode thermal evaporation source (a) side/front view (b) top view showing shielding and crucible configuration



4.2.1.1.2 Commercial source

The organic deposition panel, SVTA-10SF-4 was purchased from SVT Associates. Mounted

on a 1 0 ” flange, it consists of four commercial low temperature open-tube effusion cells, with

either AI2O3 or BN crucibles, for organic molecule evaporation. Each source is independently

controlled, with individual pneumatic shutters. The panel is situated at a 55° angle to the

substrate normal. As a result, the organic film has non-uniform thickness across the substrate

surface. The gradient between the centre and the edges of the sample may be significant

enough to affect the experimental results for films thicker than 300A.

4.2.1.1.3 Sample preparation and deposition in OMA C

Prior to deposition, the substrate surfaces were cleaned by Ar+ ion sputtering in UHV in the

analysis chamber to eliminate surface contamination by C and O. The only exception was for

organic deposition onto previously deposited thin films o f LiF or for LiF deposition onto

organic films. In those cases, the LiF and organic deposition occurred sequentially without

breaking vacuum during analysis. Once the substrates were prepared, they were transferred

into the CMAC and held in vacuum (-1x1 O' 9 Torr) until deposition occurred. The substrate

was held at room temperature throughout the deposition experiments.

For organic deposition, powders were thermally evaporated from AI2O3 crucibles in

the Knudsen cell sources onto the prepared substrates. To eliminate residual water and

oxygen, the molecules were pre-heated for at least one hour at an intermediate temperature

(0.6 Tdeposition) prior to deposition. Chamber pressure was maintained at ~5xlO ' 9 Torr. The

temperature of the source was kept fairly low, for an average deposition rate o f lA/min

(mean for experiments is 0.99±0.4 A/min). For LiF deposition, crystals were heated in an



AI2O3 crucible in the cathode source. The applied current was kept at ~20A to achieve a

deposition rate of approximately 3 A/hr.

4.2.1.2 F-M AC

All sample analysis is performed in the O-MAC, attached to the CMAC. It consists o f a Phi

ESCA 5500 multi-technique system capable of performing XPS, Auger electron

spectroscopy, scanning electron microscopy, and ion beam sputtering. The base pressure is

typically ~ l x l 0 ‘9 Torr. X-ray spectra were generated with either Mg K« (1253.6 eV)

radiation in a 54.7° geometry or with monochromated Al Ka (1486.7 eV) radiation in a 90°

geometry. The photoelectrons were analyzed by a hemispherical analyzer using 23.35eV pass

energy, with a nominal analysis area o f 800pm2 and sampling depth <50A. For angle

resolved analysis, the sample was tilted to vary the photoelectron take-off angle in the range

o f 25° and 85°. In order to examine the entire composition and structure of the interface

layer, Ar+ ion depth profiling was performed. The sputter rate was approximately 6 A/min,

calibrated for SiC^/Si structures, with a 3 keV Ar+ beam at 60° incidence angle. Typical X-

ray flux is in the range o f 1 0 10-1 0 n photons/s.

4.2.1.2.1 XPS data analysis

For most collected XPS spectra, a least-squares curve fitting analysis was carried out using

PHI MultiPak 6.1 A. The standard fitting procedure uses mixed Gaussian-Lorentizian line

shapes (Voigt summation profiles [15] as described in Appendix C) with an iterated Shirley

background (Appendix C), to account for scattering of low energy electrons. Metals and

conducting solids often deviate from purely symmetric Gaussian-Lorentzian shapes used in



Voigt profiles due to screening o f the core-hole after photoelectron excitation [16]. The

degree o f this asymmetry is proportional to the density o f states at the Fermi energy.

Insulators and semiconductors, therefore, are almost always fully defined by symmetric

Gauss-Lorentz curves. Where necessary, however, the curve fitting analysis was modified to

account for the high binding energy tail from the small energy electron-hole excitations

around the Fermi energy. In MultiPak, this is performed using a summation of the standard

Voigt formula with an exponential tail description, using an asymmetry factor. These curve

fitting formula are summarized in Appendix C.

For determination o f the core level positions, different features were used for

alignment depending on the circumstance. Most samples that were produced ex-situ to the O-

MAC were aligned based on adventitious C at 284.8 eV [17], unless otherwise noted. When

there was no adventitious C, such as during in-situ depositions, the substrate was used as a

reference point. For device peel-off, the spectra were usually aligned to the Al 2p core level

for Alq3, at 74.4 eV, if it was available. This value was determined experimentally as

described in chapter 5 and Appendix C.

For thickness and structure estimations from changes in peak intensity, the effective

attenuation length (EAL) have been calculated taking into account both the kinetic energy o f

the electron and the medium through which it will be traveling, using the National Institute o f

Standards and Technology EAL Database [18], This database allows the calculation o f any

EAL, given an inelastic mean free path (IMFP) and a transport free mean path (TMFP) for

materials o f known band gap and density. The IMPF is the average distance which an

electron will travel along a straight line path between inelastic collisions. It was generally

calculated with the Tanuma-Powell-Penn (TPP-2M) theoretically derived inelastic mean free



path equation [19] from the National Institute o f Standards and Technology Electron IMFP

Database [20]. The TMFP accounts elastic scattering, being defined as the average distance

an electron must travel before its momentum is reduced through elastic scattering alone. It

was calculated by applying the transport approximation [2 1 ] for transport cross sections

determined from the National Institute of Standards and Technology Electron Elastic-

Scattering Cross-Sections Database [22], Further details and the standard formula used are

outlined in Appendix D.

4.2.2 Cluster tool

The Kurt J. Lesker OLED cluster tool (figure 4-4), assembled in 2003, is very similar to the

MAC system, with multiple deposition chambers around a central distribution chamber. This

is a dedicated deposition system, with no in-situ analysis capabilities. The cluster tools

include a central distribution chamber, a loadlock chamber, a plasma treatment chamber, a

sputtering chamber, an organic deposition chamber, and a metallization chamber. The base

opressure is maintained at 10" Torr. As the cluster tool is designed for large substrates, the

sample stage is rotated during deposition to ensure uniform film thickness across the surface.

Deposition rates in this system are approximately lA/s for the organics, and slightly slower

for inorganics. Typical deposition rates for LiF range between 0.2-0.5A/s, and for metals

around from 0.6-5A/s. Thicknesses were determined with an Inficon XTM/2 quartz crystal

microbalance as in the OMAC, calibrated by X-ray photoelectron spectroscopy, mass and

profilometry.

Any sample that is referred to as a “device structure”, both for electrical and chemical

analysis, is deposited in the cluster tool on 2”x 2” glass substrates. After the substrates are

treated by oxygen plasma for 10 min in the plasma chamber, they are transferred to the



sputtering chamber where an ITO film is deposited by RF sputtering at a power of 45 W and

an argon pressure of 8.5 mTorr. First, a grid shadow mask was used to define the ITO anode

structures (ITO sheet resistance was -300 Q/sq), then organic molecules, and cathode

materials were sequentially deposited by thermal evaporation in the organic and metallization

chambers.

For samples made in the cluster tool, analysis was always performed completely ex-

situ, with varying amounts o f air exposure. Generally, samples were analyzed within 20 mins

of breaking vacuum after deposition, or were kept in vacuum after minimal air exposure until

analysis. The only exceptions were the samples used in the oxidation rate and shelf time

studies, which were deliberately exposed to the ambient environment in the laboratory.

Figure 4-4 Kurt J. Lesker OLED cluster tool in the clean room


Chapter 4 Experimental

4.3 In-situ peel off method

76

Attached to the <D-MAC is the delamination chamber, used for in-situ peel-off experiments,

designed to expose buried interfaces for study by XPS. The chamber consists of a loading

fork to allow transport into the XPS chamber, a mechanical armature used to apply a lifting

force to the tape for peeling, and an elevated window to allow observation o f the sample

surface during the peel-off. The peel-off procedure, performed at <lxlO ' 6 Torr, is described

below.

The conventional method to characterize the buried interface is using ion sputtering to

remove the top layer. For conventional inorganic materials such as SiCVSi interface [23,24],

this works reasonably well. A problem with sputter profiling for organic molecules, however,

is the significant interface mixing that may occur due to ion beam bombardment. With

artefacts introduced from sputtering, often no meaningful data can be derived about the

interfacial structure [25].

The well known tape peel test for adhesion strength [26] at thin film interfaces has

been previously applied in ex-situ degradation studies of OLEDs [27,28]. Although Murase

et al. [29] have made use o f the technique for peel-off in air, our facility allows peeling to be

performed under vacuum in the delamination chamber.

The peel-off procedure is performed by applying a conductive and adhesive tape on

the sample surface. A tensile force is then applied along the surface normal to pull the tape

o ff All of the cathode metal was found to adhere to the tape while the rest o f the organic

layers adhere to the substrate, and thus the peel-off leads to almost perfect cleavage at the

organic/metal interface. Figure 4-5 shows a schematic o f the peeling procedure for exposure

o f both the cathode and organic sides of the interface.


Reproduced

with perm

ission of the

copyright ow

ner. Further

reproduction prohibited

without

permission.

Side viewsam ple configuration during p eel-o ff

I . glass substrate I

cathode

Top View

^ applied

\carbonjape

attach \ inverted j

sample giass substrate

substrate holderorganic

/cathode (analects, region)

peel-off

(back side) cathode

cathode

.glass.suhstratel]

indicates cleavage region during peel-o ff

carbi

glassx substrate

(front side)

cathd't

applied

attachtape

pee I-off •

scotch tape

substrate holder

organic underlyingcathode

(analysis region)

Figure 4-5 Schematic of peel-off procedure. Top panel shows the removal of the glass substrate to expose the cathode surface for analysis. Bottom panel shows the removal of the cathode layer to expose the organic surface for analysis indicated by the circled area. Conductive carbon tape was used in both instances to adhere the sample to the sample holder to minimize any charging effects. The cleavage plane during peel-off is indicated by the heavy dotted line in the “side view” section.

<1o

Chapter 4 Experim

ental


After being peeled off, the buried interface becomes two surfaces; the organic film

surface and buried cathode surface. Depending on the positioning o f the tape on the sample

surface, either the organic surface or the cathode surface may be independently analyzed;

however, it is not possible to perform analysis on both sides of the interface simultaneously.

For the examination o f the cathode layer, the sample was placed cathode side down on

electrically conductive carbon tape on the sample holder. The glass substrate and organic

layers were then removed using adhesive tape as shown in the top panel o f figure 4-5,

leaving behind the metallic film for analysis. For the organic layer, the cathode was removed

from a different sample cut from the same glass “wafer”.

To ensure that the samples were not damaged during scoring and cutting o f the

individual samples with a diamond cutting tool, the wafer substrate was mounted in a home

built holder, as shown in figure 4-6 below.

Wafer substrate

Front view

Top view

Figure 4-6 Sample holder schematic for substrate scoring before peel-off

The adhesive tape was positioned on the sample outside the vacuum environment,

and the entire sample was loaded into the delamination chamber for in-situ peel-off. Once the

chamber was evacuated to a high vacuum condition (~ 10 6 Torr) by a turbomolecular pump,

the surface of interest was exposed. Subsequently, the substrate holder was loaded into the

analysis chamber for characterization. This method showed a high degree of reproducibility,



with cleavage in the top organic layer within ~40A of the interface. For example, four

different peel-offs o f the same metal/organic interface yields the same XPS results. This

method is also applicable to a wide variety o f organic molecules and cathode materials,

consistently indicating cleavage at the cathode/organic interfaces.

4.4 Other analysis techniques

Some high resolution scanning electron microscopy was performed at the Institute for

Microstructural Studies at the National Research Council in Ottawa on a Hitachi S4700. As

well, atomic force microscopy was performed on some sample surfaces using a Digital

Nanoscope E with a 15 pm scanner in the Department o f Materials Science and Engineering

at the University of Toronto. Electrical luminance-current-voltage (L-I-V) measurements of

device structures were performed using an HP 4140B pA meter coupled with a Minolta LS-

110 luminance meter in the Lu group clean room at the University o f Toronto.

4.5 References1 John E. Mahan, Physical Vapour Deposition o f Thin Films (John Wiley and Sons, New York, 2000).

2 S. M. Rossnagel, J. Vac. Sci. Technol. A 21, 574 (2003).

3 Milton Ohring, Materials Science o f Thin Films, 2nd edition (Academic Press, San Diego,2002), Chap. 3.

4 S. G. Brush, The Kinetic Theory o f Gases: An anthology o f classic papers with historical commentary (Imperial College Press, London, 2003).

5 L. F. Cheng, L. S. Liao, W. Y. Lai, X. H. Sun, N. B. Wong, C. S. Lee, and S. T. Lee,Chem. Phys. Lett. 319, 418 (2000).

6 S. Mochzuki, M. Sasaki, and R. Ruppin, J. Phys.: Condens. Matter 10, 2347 (1998).

7 K. Tanigaki, S. Kuroshima, and T. W. Ebbesen, Thin Solid Films 257, 154 (1995).

8 M. F. Butman, A. A. Smirnov, L. S. Kudin, and Z. A. Munir, J. Mater. Synth. Process. 8 ,93 (2000).



9 G. M. Rothberg, M. Eisenstadt, and P. Kusch, J. Chem. Phys. 30, 517 (1959).

10 M. Eisenstadt, J. Chem. Phys. 29, 797 (1958).

11 M. Knudsen, Annal. Physik IV 48, 1113 (1915).

12 I. Langmuir, J. Am. Chem. Soc. 38, 2221 (1916).

13 E. B. Halac, M. Reinoso, A. G. Dall'Asen, and E. Burgos, Phys. Rev. B 71, 115431 (2005).


15 P. M. A. Sherwood, in Practical Surface Analysis, 2nd edition, edited by D. Briggs and M. P. Seah (Wiley & Sons Ltd, New York, 1990), Vol. 1, App. 3, p.573; (b.) W. Voigt.Munch. Ber. 1912, 603 (1912).

16 S. Doniach and M. Sunjic, J. Phys. C 3, 285 (1970).

17 J. F. Moulder, W. F. Stickle, P.E. Sobol, K.D. Bomben, Handbook o f X-ray Photoelectron Spectroscopy, edited by J. Chastain, and R.C. King, Jr. (Physical Electronics Inc., Eden Park, MN, 1995).

18 C. J. Powell and A. Jablonski, NIST Electron Effective-Attenuation-Length Database - Version 1.0, National Institute o f Standards and Technology, Gaithersburg, MD, (2001).

19 S. Tanuma, C. J. Powell, and D. R. Perm, Surf. Interface Anal. 21, 165 (1994).

20 C. J. Powell and A. Jablonski, NIST Electron Inelastic-Mean-Free-Path Database - Version 1.1, National Institute o f Standards and Technology, Gaithersburg, MD, (2000).

21 A. Jablonski and S. Tougaard, J. Vac. Sci. Technol. A 8 , 106 (1990).00 A. Jablonski, F. Salvat and C. J. Powell, NIST Electron Elastic-Scattering Cross-Section

Database - Version 3.0, National Institute o f Standards and Technology, Gaithersburg, MD, (2002).

23 See, for example, L.C. Feldman and J.W. Mayer, "Fundamentals of Surface and Thin Films" (North-Holland, Amsterdam, 1986).

24 Z.H. Lu and D. Grozea, Appl. Phys. Lett. 80, 255 (2002).

25 P.R. Norton, private communication.0 f \ M. Ohring, The Materials Science o f Thin Films (Academic, Toronto, 1992), p. 25.

27Y. Sato and H. Kanai, Molecule. Cryst. Liq. Cryst. 253, 143 (1994).


29 A. Murase, M. Ishii, S. Tokito, and Y. Taga, Anal. Chem. 73, 2245 (2001).


Chapter 5

Metal/Alq3 interface structures1

5.1 Introduction

Since the first practical organic light-emitting diode (OLED) was reported [1], there has been

widespread interest in understanding and improving OLEDs, in particular metal/organic

interfaces. Most organic materials, even the predominantly used Alq3, have relatively poor

electron transport characteristics; therefore, the metal/organic interfaces are known [1 ,2 ,3,4]

to play a vital role in device performance such as turn-on voltage, reliability, and quantum

efficiency. Metal/organic interface formation is rather complex as it may involve molecular

fragmentation and atomic reaction/diffusion, as opposed to conventional metal/semi-

conductor interfaces where the interface formation follows well-established material thermo

dynamics. A majority o f OLED structures are produced by layer-by-layer deposition from

1 First appeared in a slightly different format as Applied Physics Letters 81(4) 766-768, Copyright 2002, American Institute of Physics (reproduced with permission).

- 81 -


Chapter 5 Metal/Alq3 interface structures 82

anode to cathode, with the final major fabrication step being the deposition of the inorganic

cathode onto an organic layer [1-4]. The deposition o f metallic cathode materials on an

organic substrate has a different reaction sequence than organic deposition onto metal

substrates [5], A wide variety o f metal/organic systems have been examined in order to

optimize OLED performance, and Mg and Mg:Ag alloy cathodes have shown great promise,

though no one system predominates. There have been several studies [6,7,8] using XPS to

track down initial reactions when a metal atom is deposited upon an organic film surface.

Previous studies o f the interface have either deposited ultra-thin layers o f cathode on

the organic, analysing each layer [5,7,9] or deposited thick layers, sputtering the metal to

expose the interface [10]. Both o f these methods are “top down” analyses from the

air/cathode interface. As the interface is buried in device structures, most photoelectron

investigations using these techniques have not been able to directly examine both the cathode

and organic side o f the interface in devices. Although useful information has been gained

from these types o f studies, the structure of buried metal/organic interfaces, especially in an

operating OLED, is not completely understood. In this chapter, our findings on the interface

structures o f various buried metal/Alq3 interfaces from working OLEDs, using XPS are

discussed. In order to establish a clear interface reaction model, four different metals were

used as the cathodes, which are Au, Ag, Mg, and Mg:Ag alloy. Mg and Mg alloys are

standard cathodes for OLEDs, and Au represents a model inert metal.

5.2 Experimental

Functional OLEDs were fabricated using a deposition procedure described in chapter 4. In

this case, all materials were deposited in the same vacuum chamber (base pressure: 5x1 O' 6



torr) from resistively heated W boats at a rate o f 3 A/s onto ozone-cleaned ITO substrate. The

samples consisted o f 1300A layer o f metal (Ag, Mg, or Au) onto 1000A of Alq3 on ITO

coated glass substrates. A similarly produced OLED with structure ITO/15A CuPc/600A

NPB/75A Alq3/2 0 0 0 A Mg9o:Agio cathode was also analyzed. The XPS spectra for all

samples were generated using Mg Kq radiation (1253.6eV) in a 54.7° geometry with pass

energy o f 29.35eV under a base pressure of 1x10 9 Torr. In order to examine the entire

composition and structure o f the interface layer, Ar+ ion depth profiling, with a sputter rate of

approximately 24A per cycle for Au, Ag and Mg and 20A per cycle for Mg:Ag, was

performed with a 3 keV Ar+ beam at 60° incidence angle.

5.3 Results and discussion

Table 5-1 summarizes the XPS data obtained on the “as-peeled-off’ surfaces. For all buried

Alq3 surfaces, the N/Al ratio is found to be 3.0, as expected from its molecular structure. On

the cathode side, we also observed N and A1 peaks. These data suggest that the peel-off

process occurred perfectly between un-reacted Alq3 and the reacted region attached to the

metal cathode. Upon further examination o f the metal side, we found that the N/Al ratios on

buried Au and Ag surfaces are about 3, while that on buried Mg and Mg:Ag alloy surfaces

are about 1.1. The results suggest that the N and A1 species on the Au and Ag surfaces are

likely in the form of Alq3 attached to the cathodes, while that N and A1 species on the Mg

and Mg: Ag cathode are likely related to new compounds formed at the interface.



Table 5-1 XPS measured N/Al ratios on various buried surfaces. The sensitivity factors are: 0.472 for N Is, 0.250 for A1 2p, and 0.333 for A1 2s, respectively. The theoretical N/Al ratio is 3, calculated based on Alq3 molecular structure. = = = = _ = _ _ =_ _ _ _ _ _ = = = = = = = ^ ^

Alq3 Surface (all samples) Au Surface3 Ag Surface Mg Surface Mg: Ag Surface

N/Alratio 3.0 2.83 3.12 1.1 1 .2

a) Residual amount of Au was detected on the Alq3 surface; A1 2s was used to determine N/Al ratio.

The conclusion is further supported by A1 core level binding energy position. Using

Au 4 /7/2 at 84.0 eV as a reference, a simple peak with a binding energy at 74.4eV has been

found for A1 2p2 on all buried Alq3 surfaces and on Au and Ag cathodes; while a complex A1

2p peak has been found on Mg and Mg:Ag cathodes. Figure 5-1(a) shows A1 2p core level

spectra recorded from Ag, Mg, and Mg:Ag cathode surfaces. Figure 5-1(b) shows that the A1

2p spectra observed on Mg and Mg:Ag cathodes may be well-fitted by two doublet peaks;

one at 72.7 eV and another at 74.4 eV. The A1 2p spin-orbit doublet was not resolved because

of insufficient instrumental resolution. For curve-fitting A1 2p, we used the Voigt summation

function, with 0.4 eV spin-orbit splitting [11] for the A1 2pm,m doublet. The peak at 72.7 eV

is similar to that obtained on a metallic aluminium surface, and therefore is attributed to the

formation of metallic aluminium species on the cathodes. The peak at 74.4 eV is attributed to

that of residual Alq3 on the Mg and Mg:Ag cathodes. This attribution is further supported by

the fact that N/Al ratio is ~ 3 when only the peak at 74.4 is used in the calculation.

The Au/Alq3 system showed the greatest difficulty for alignment on analysis o f the

cathode side of the interface, as the strong signal from the Au 5pm core level, expected at 74

eV, obscures the A1 2p core level signal3. On the Ag cathode, where both A1 2s and A1 2p

2 See Appendix C for details on alignment and determination of A12p3/2 position.3 Sputter depth profile measurements into the Au cathode confirm that the signal observed is a result o f the cathode and not due to Al.



core levels were visible, the behaviour o f the A1 2s level was consistent with the expected

value for Alq3. Therefore, to examine A1 on the cathodes side for Au, the A1 2s core level

was observed instead.

Figure 5-1 Top panel shows A1 2p core level spectra recorded on various as peeled off metal surfaces: Ag surface shown as circles, Mg surface shown as open triangles and Mg:Ag alloy surface shown as solid circles. The bottom panel shows curve fitting results of A1 2p recorded on the Mg:Ag surface. The experimental data (solid circles) can be well fitted by the sum (solid line) of two separate spin-orbit doublet peaks (dashed lines), one metallic state at 72.7 eV and another Al3+ state at 74.4 eV.

78 76 74 72 70 68

Binding Energy (eV)

In order to establish a depth distribution o f various species on the buried metal

surfaces, we used 3 keV Ar+ ion bombardment to carry out XPS depth profiles. The sputter

rate is calibrated for SiCVSi structures. It should be pointed out that sputter rates vary

depending on the material systems. For all metals studied here, they generally have a higher

sputter rate than SiC>2, so the depth o f the reaction zone is only relatively comparable. Figure

5-2 compares A1 2p core levels profiled to various depths on the metal cathodes. For Au, the

results with A1 2s were similar to that o f the A1 2p profile on the Ag cathode. All spectra are

shown with intensities as recorded. The figures show that A1 2p peak is very weak on Ag

Al 2 p (a)

c3_QL.3*

c

C athode o Ag* Mg* Mg:Ag

* CO*p ° * o o

-4—*cAl 2 p■a

0NcoEoz



cathode at the as-peeled-off surface, i.e. d=0 A. With sputtering, this feature and the N Is core

levels (not shown here) were no longer visible, suggesting that the organic residue on the

cathode surface from the peel-off was completely removed during sputtering. The Al 2p core

levels data on Mg and Mg:Ag cathodes, however, exhibit a rather remarkable evolution.

There are two types o f aluminium on the as peeled off surface (d=0 A), as discussed above.

The XPS profiles show that the Al3+ species decrease in intensity while the metallic Al

species increase in intensity with increased depth. The data suggest that significant Al

diffusion into the Mg and Mg:Ag cathodes have occurred.

1 4 4A Figure 5-2 XPS depthprofile of as-recorded Al

96A 2p core levels obtainedfrom: (a) Ag cathode, (b) Mg cathode and (c) Mg:Ag alloy cathode.

(c)

80 78 76 74 72 70 6 8

Binding Energy (eV)



The Mg species also showed an interesting evolution through the thickness of the

cathode, as shown in figure 5-3 for the Mg 2p core level depth profile obtained on the peeled

off Mg and Mg:Ag cathode surfaces. At the as peeled-off surface (d=0A), the Mg 2p core

level was found to be centred at ~50eV for both cases, indicating the existence o f metallic

Mg at the interface. A second peak at a higher energy after sputter removal o f ~ 50A was

found on the cathodes. In both cases, a higher binding energy state was observed, indicating

the presence o f Mg oxides. It is very important to note that this layer of Mg oxides is

distributed away from the Alq3/metal interfaces, with a metallic Mg sandwiched in between.

Upon further sputtering, the Mg 2p evolution from metallic state to oxidation state is rather

clear and abrupt on Mg cathodes. Whereas, for Mg:Ag cathodes, the asymmetry o f the Mg 2p

core level suggests a rather mixed state persists on Mg:Ag cathodes throughout the depth

profile. With further sputtering into the cathode, the intensity o f the oxide peak decreased and

the metallic peak reappeared. The data suggests a well-defined Mg/MgO/Mg/Alq3 layer-by-

layer structure formed at the Mg/Alq3 cathode interface, while a rather mixed phase is visible

at the Ag:Mg/Alq3 interface. The pure Ag cathode shows no reaction to the Alq3, with the

binding energy of the 3d core level unchanged through the thickness o f the cathode layer.

The Au cathode shows a similar result, with no reaction to the Alq3 and consistent binding

energy values through the thickness of the cathode layer4.

4 Due to the interference from the Au cathode, the charging effects could not be accounted for using the Al 2p as a reference.



d = OA

"Dd).NTOE!_o

Z

Mg 2p v (b);

192 A :

X v_ _ ^ J4 4 A ■

V :

\ 48A

V _ d = OA :

Figure 5-3 XPS depth profile of intensity normalized Mg 2p core levels obtained from: (a) Mgcathode and (b) Mg:Ag alloy cathode.

54 53 52 51 50 49 48 47 46

Binding Energy (eV)

In Figure 5-4, the interfacial chemical structure inferred from the XPS data is summarized

using schematic diagrams for the four different metal cathodes.

MgOx,Mg, Ag, Al MgOx, Al Ma. Al

Au

Figure 5-4 Schematic summary of various interface structures: (a) Mg:Ag/AIq3 interface, (b) Mg/Alq3 interface, (c) Ag/Alq3 interface, and (d) Au/Alq3 interface.



Previous studies o f the cathode-organic interface for a variety o f metals indicate that

the inorganic cathodes tend to diffuse into the organic layer during deposition [5,7,10,12,13,

14] most likely due to hot metal atom impingement effects on the weakly bonded organic

surface. This impingement driven diffusion effect was observed on the organic side o f the

interface, for all cathodes examined in this investigation. The XPS sputter depth profiles for

the Au and Ag cathodes indicate no chemical interaction between the organic and inorganic

layers. This is expected for a diffusion coupled interface structure, with two independent

layers of cathode and Alq3 on either side of the diffusion layer (Fig 4. (c) and (d)). The (Au,

Ag)/Alq3 interface formation, therefore, is rather simple; no interface chemical reaction

occurs. The (Mg, Mg:Ag)/Alq3 interface formation, however, follows rather complex

reaction/diffusion processes.

Earlier XPS studies have indicated that the structure o f Alq3 [5-7] is modified by Mg.

The existence o f oxide and diffusion layers within the cathode beyond the persistence o f N,

suggests that Mg is scavenging the O from the phenoxide ring in the quinolate ligand, leading

to molecular fragmentation, rather than simple Mg attachment as has been previously

proposed [5,7,15,16,17]. Since both the initial deposition and the peel-off were performed

under high-vacuum conditions, Mg oxidation observed here would most likely occur through

the liberation o f oxygen from the organic layer. The shift in the Mg 2p core level (figure 5-2)

indicates that the reaction zone within the cathode is substantial. A previous study on the

Mg:Ag cathode has also shown the existence of such a reaction zone [18]. The high reactivity

o f Mg would allow for the interaction of up to three Mg atoms with each Alq3 molecule, as

all the ligands are symmetrical. The formation of Mg oxides would likely occur at the

expense of the A l-0 bonds in the molecule, since the O is the most thermodynamically



favourable site o f interaction [16-19]. With a sufficient supply of Mg atoms, the molecule

could completely fragment along the phenoxide bonds, liberating Al from the metal chelate,

following a reaction:

M g + Alq^ —> M g + A l + Qx (5-1)where Qx stands for fragmented hydroxyquinoline.

Although the chemical nature of Qx is difficult to quantify, XPS data shown in Table

5-1 indicate N deficiency. This suggests that some component o f Qx are gaseous N species

which have evaporated either during initial stage o f deposition or after the interface being

peeled open in the vacuum.

i IBased on the fact that the chemical state o f Al in Alq3 is Al ; the above reaction may

be rewritten as

2M g + (y ,)A > A -> 2M gO + {y,)AI (5-2)

This is a well-known oxidation-reduction reaction and is thermodynamically

favoured. At room temperature, for example, the reaction will lead to a reduction of ~ 84 kJ

o f free energy [20].

Related to this oxidation are diffusion processes. For both Mg/ and Mg:Ag/Alq3

interfaces, XPS data show metallic Mg species and intact Alq3 molecules. This indicates that

the oxidation-reduction reaction is limited by the reaction rate, rather than limited by Mg

diffusion into the interface. It is possible that diffusion will be significantly slowed down

when the reacted region becomes thick enough to act as an effective diffusion barrier. It

implies that the thickness o f interface oxides may grow with time, and thus the OLED

driving voltage may grow as a function of time as well. Should oxide growth proceed with an

island type of growth pattern, dark spot formation will be a likely consequence.



For the alloy cathode, the presence o f Ag in the alloy would generally be expected to

enhance Mg diffusion by opening up more vacancy pathways due to large Ag atom

substitution into a smaller Mg lattice [21]. However, an ultrathin layer o f Ag between the

organic and cathode has been previously used to block Mg migration [22], Mg diffusion

could, therefore, be inhibited due to Ag atoms, perhaps through the formation o f (3-MgAg

[23], However, this limited diffusivity has little impact on the molecular breakdown at the

interface, which occurs in either case.

The second diffusion species is the metallic Al. Figure 5-2 shows that Al diffusion is

very extensive. The Al 2 p core levels are very strong even when ~ 200A of cathode material

has been removed. In general, the diffusion o f metallic Al into Mg is expected because o f a

high solid solubility of Al in Mg [23]. It is also a known fact [24], however, that Al/Mg thin

film systems do not undergo a reaction at room temperatures. The unusually high diffusion

observed here may be attributed to the fact o f Mg diffusion to the organic surface, driven by

the formation o f more stable MgO. The Mg diffusion would generate abundant vacancies

VMg that would in turn serve as pathways for metallic Al diffusion.

The limited reaction zone for the alloy cathode could, therefore, be related to the

suppressed diffusion of Mg. Though no room temperature diffusion data exist, extrapolating

from high temperature data indicates that the diffusion constant would be four orders of

magnitude less for Al in Ag (~lxl0~29 cm2/s[25]) than in Mg (~8xl0‘25 cm2/s[26]). The

presence of Ag, by slowing down both Mg and Al diffusion, limits the extent o f diffusion of

Al visible by XPS compared to that observed with the Mg cathode alone.



The Ag and Au cathodes show no such diffusion of Al into the cathode. In the

presence o f metallic Al, both Au and Ag would be expected to show strong diffusion.

Metallic Al shows a strong solid solubility in Ag even at low temperatures, indicating a high

driving force for diffusion [23]. The Au-Al thin film system, with excess concentration of

Au, has a high driving force to form a stoichiometric compound, AU2AI [24]; therefore, if

metallic Al were present, there would be a high driving force for reaction with Au, which has

not been observed. Since Al 2p and N Is core levels disappear upon sputtering, only intact

molecules exist at the interface. This implies that the Al liberating interfacial reaction

observed in the Mg case was not occurring upon deposition of less reactive cathode

materials.

This fragmentation behaviour may not have been observed by other researchers due

to the reaction zone thickness. Estimated to be >200-300A (see both figures 5-2 and 5-3), this

zone is well beyond the detection depth limit o f traditional overlayer techniques. Mg may,

like Ca [27], have a critical thickness for reaction initiation that could not be observed

without an examination of the buried interface. This could suggest that at the earliest stages

of interface formation, anion formation occurs with N modification through charge transfer

from the metal. Upon further deposition, however, Mg oxidizes and fragments the molecule.

This could also explain why the N Is shoulder observed by other researchers during in-situ

deposition was not observed in this work. Alternatively, the N Is shoulder could be explained

by the weakly bonded N2 gas from the fragmented quinolate trapped by the overlayers during

deposition.


Chapter 5 Metal/Alq3 interface structures

5.4 Summary

93

The new information on the contact formation process, gained by the study of buried

interfaces in conventionally fabricated devices using a unique peel-off technique, allows a

connection to be drawn between the equilibrium structures and the non-equilibrium

structures observed during traditional investigations o f built-up interfaces by monolayer

deposition. The interface structures found in Mg:Ag, Mg, Ag, Au/Alq3 systems show that

cathode metals can be broken up into two broad categories, those that form an interfacial

reaction layer and those that do not. Mg and Mg:Ag cathodes have long reaction zones, with

complicated oxidation/diffusion processes occurring at the interface. Since the centralized

metal atom, Al, in Alq3 is in a 3+ oxidation state, it is expected to behave approximately like

AI2O3. Using this approximation, the likelihood of this reaction can be predicted by

modelling the organic reaction as an inorganic oxidation-reduction metal exchange reaction.

5.5 References


2 Y.F. Yiew, H. Aziz, N.X. Hu, H. Chan, G. Xu, and Z.D. Popovic, Appl. Phys. Lett. 77, 2650 (2000).

3 M. Ikai, S. Tokito, Y. Sakamoto, T. Suzuki, and Y. Taga, Appl. Phys. Lett. 79, 156 (2001).

4 For a recent review, see for example, I.H. Campbell and D.L. Smith, in Solid State Physics, edited by H. Ehrenreich and F. Spaepen (Academic Press, New York, 2001), Vol. 55, p .l.


6 N. Isomura, T. Mitsuuoka, T. Ohwaki, and Y. Taga, Jpn. J Appl. Phys. 39, L312 (2000).

7 P. He, F.C.K. Au, Y.M. Wang, L.F. Cheng, C.S. Lee, and S.T. Lee, Appl. Phys. Lett. 76, 1422 (2000).

8 C. Shen, I.G. Hill, A. Kahn, and J. Schwartz, J. Am. Chem. Soc. 122, 5391 (2000).

9 Z. H. Ma, S. L. Lim, K. L. Tan, S. Li, and E. T. Kang, J. Mater. Sci. - Mater. El. 11, 311(2000).



10 W. Song, S. K. So, J. Moulder, Y. Qiu, Y. Zhu, L. Cao, Surf. Interface Anal. 32, 70 (2001).

11 S. Hufiier, Photoelectron Spectroscopy (Sringer-Verlag, Berlin, 1995), p.456.

12 Y. T. Tao, E. Balasubramaniam, A. Danel, B. Jarosz, P. Tomasik, Chem. Mater. 13, 1207(2001).

13 G. Gu, G. Parthasarathy, P. E. Burrows, P. Tian, I.G. Hill, A. Kahn, and S.R. Forrest J.Appl. Phys. 86,4076(1999).

14 L. Zou, V . Sawate'ev, J. Booher, C. H. Kim, J. Shinar, Appl. Phys. Lett. 79, 2282 (2001).

15M.E Thompson, P.E. Burrows, and S.R. Forrest. Curr. Opinion Solid State Mater. Sci. 4,369 (1999).

16M. G. Mason, C. W. Tang, L-S. Hung, P. Raychaudhuri, J. Madathil, D. J. Giesen, L. Yan,Q. T. Le, Y. Gao, S-T. Lee, L. S. Liao, L. F. Cheng, W. R. Salaneck, D. A. dos Santos, and J. L. Bredas, J. Appl. Phys. 89, 2756 (2001).

17 R. Q. Zhang, W. C. Lu, C. S. Lee, L. S. Hung, and S. T. Lee, J. Chem. Phys. 116, 8827 (2002). b. R. Q. Zhang, X. Y. Hou, S. T. Lee, Appl. Phys. Lett. 74, 1612 (1999).

18 X.D. Feng, D. Grozea, A. Turak, Z.H. Lu, H. Aziz, and A.-M. Hor, MRS Symp. Proc., Organic and Polymeric Materials and Devices - Optical, Electrical, and Optoelectronic Properties, San Francisco, v. 725, P.4.8.1 (2002).

19 A. Curioni and W. Andreoni, J. Am. Chem. Soc. 121, 8216 (1999).90 M. Ohring, The Materials Science o f Thin Films (Academic Press, Toronto, 1992), p.25;

Thermochemical Data o f Pure Substances, 3rd edition, edited by I. Barin (VCH Publishers, New York, 1989), Vol. 1

21 Diffusion Phenomena in Thin Films and Microelectronic Materials, edited by D. Gupta, and P.S. Chan (Noyes Publications, Park Ridge, NJ, 1988)

22 M. Kiy, I. Gamboni, U. Suhner, I. Biaggio, and P. Gunter, Synth. Met. 111-112, 307 (2000).93 See, for example, Binary Alloy Phase Diagrams, edited by T.B. Massalski (ASM, Metals

Park, Ohio, 1986).

24 V. Simic and Z. Marinkovic, J. Mater. Sci. 33, 561 (1998).25 Smithells Metals Reference Book, 7th edition, edited by E. A. Brandes, and G. B. Brook

(Butterworth-Heinemann, Oxford, 1998), Chap. 13

26 G. Moreau, J. A. Comet, and D. Calais, J. Nucl. Mater. 38, 197 (1971).

27 V. Choong, Y. Park, Y. Gao, T. Wehrmeister, K. Mullen, B. R. Hsieh, and C. W. Tang, Appl. Phys. Lett. 69, 1492 (1996).


Chapter 6

LiF/metal bilayer structures I - Case of Al/LiF

6.1 Introduction

Organic optoelectronics have garnered considerable interest over the last 30 years for the

enhanced optical flexibility offered by the use o f organic semiconductors. With organic

semiconductor devices, the electron injection characteristics o f the cathode/organic interface

play a critical role in overall device performance. Early on, low work function metals

[1,2,3,4,5] were used as cathodes due to the presumed low interfacial barrier to electron

injection. However, such metals also exhibit high chemical reactivity with both the ambient

environment and the organic active layers themselves [3,6,7], leading to inefficient

performance and short device lifetimes. Significant improvements to both performance and

lifetime have been observed with multilayered cathode structures. A thin, 5-10A, ionic

-95 -


Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 96

insulating interlayer, such as LiF, between the metal cathode and the organic layer has been

particularly successful in improving the injection properties across the interface [8]. As many

cathode metals strongly interact with the organic layers, one of the proposed mechanisms for

performance improvement with LiF interlayers has been the suppression of interfacial

breakdown reactions between the reactive metals and the susceptible organic active layers

[9]. In order to understand the potential interfacial reactions with such multilayer structures

in the device, it is important to first understand the oxidation behaviour of significant metal-

interlayer combinations under ambient conditions. Most o f these multilayer cathodes use

standard industrial materials such as Al and Mg:Ag alloys as a primary cathode component.

As Al also has well described oxidation behaviour, LiF coated Al surfaces represent ideal

systems for a study o f multilayer cathode oxidation.

Much research has been carried out on Al oxidation [10,11,12,13,14,15], especially

regarding the sensitivity o f the oxidation processes to small surface activity changes.

Deliberate surface activity modification with overlayer coatings is widely utilized, for both

passivation and activation of metal surfaces [16]. On Al, which self passivates at room

temperature, overlayers generally decrease the oxidation rate by forming a physical barrier

on the surface. G. Hass and collegues showed considerable lifetime improvements for LiF

coated Al mirrors [17,18,19], through total blocking o f the Al surface with LiF overlayers

thicker than 150A [18]. With extremely thin, 5-10A, layers o f LiF as used in organic

optoelectronics, the probability of a complete interlayer blocking the surface is unlikely, and

the impact o f such thin overlayers on the oxidation kinetics has been largely unexplored. This

is of great interest for interfacial device structures as any change in the oxidation kinetics

during surface exposure to air would be magnified within a device, where oxygen can only

come from disruption of the molecule or from lateral diffusion.



Recent results on organic light emitting diodes with C6o layers indicate that LiF/Al

cathodes prolong the device lifetime, presumably due to oxidation prevention at the

organic/cathode interface [20]. In this chapter, to clarify the effect o f LiF at these interfaces,

we discuss the use o f XPS to investigate the oxidation kinetics and by-products for Al

surfaces coated with thin layers of LiF, and for devices with Al/LiF cathodes. The observed

passivation of Al surfaces suggests that lifetime improvements in devices can be tailored by

controlling the cathode surface activity with “lattice matching” interlayers.

6.2 Experimental

6.2.1 Sample preparation and analysis

The coated metal structures for surface oxidation were produced using the Kurt J. Lesker

OLED cluster tool by thermal evaporation of 5000A of Al or Mg onto Si (100) substrates

under 10'6 Torr vacuum. Shadow masks were then used to deposit thin layers o f LiF on half

the surface at a rate of 0.5A/s. Samples nominally coated with 5 or 10A of LiF and uncoated

metal surfaces from the same wafer substrate were then exposed ex-situ to laboratory air

(300K, 20-30% relative humidity) for various times ranging from 20 mins to 3000hrs.

Functional OLEDs were fabricated using the deposition procedure described in

Chapter 4. In this case, all layers were deposited in OLED cluster tool to give a number of

different device structures. To examine the longevity of devices, structures o f

glass/ITO/600A NPB/200A C6o/xA LiF/2000A Al were deposited, with x varying between 5-

30A. A sample with a 100A LiF layer was also fabricated. C6o was chosen as the primary

electron transport layer for this study for two reasons - first, to use a system where interfacial


Chapter 6 LiF/metal bilayer I: Case of Al/LiF 98

oxidation is thought to be a major cause of device failure1; and second, to simplify the

analysis of the Al 2p core level during peel-off by eliminating the contribution from the Alq3

molecule. As this second reason is less important as the LiF thickness increases, device

structures similar to those o f the C60 based devices were deposited with an Alq3 electron

transport layer with LiF interlayer thicknesses of 100 and 200A.

Some of these device structures were peeled-off in-situ and examined by the XPS

using monochromated Al Ka (1486.7eV). The device performance for some devices

deposited on the same glass substrate was also measured. For seven days, the devices were

electrically stressed in air to a maximum voltage of 5V. After every test, the samples were

left exposed to laboratory air overnight (295K, 25% humidity).

Spectra were generated using a monochromated Al Ka (1486.7eV) source with a

23.35eV pass energy. The photoelectron take-off angle was varied between 25° and 85° for

angle resolved analysis o f the surface. All spectra were aligned based on adventitious C at

284.8eV [21], unless otherwise noted. Least-squares curve fitting analysis was carried out as

described in chapter 4. The shape and area of the metallic core level was kept constant

during curve fitting for all exposure times, as listed in table 6-2 in section 6.3 below.

6.2.2 Thickness and coverage determination

XPS is particularly suited for surface oxidation studies. By monitoring the subtle changes in

the chemical state o f near-surface atoms, the chemical activity o f the surface can be

determined. In addition, angle resolved XPS can be used to reliably determine the overlayer

thicknesses [22]. In this way, the oxidation kinetics can be observed concurrently with the

oxidation chemistry, giving useful information about the overall oxidation process.

1 See section 2.1.3.2 and references therein, as well as section 6.4 in this chapter.



The intensity ratios o f various photoelectron peaks were determined from the curve

fitting analysis. The oxide thickness was then estimated from the ratio o f the intensity o f the

oxide and metallic components of the metal core level by using various overlayer models, as

described in section 3.5.

For simple metal surfaces, assuming a uniform oxide film forms on the surface

(uniform overlayer model [22]), the oxide thickness can be estimated from equation 3-11

applied to the metal and oxide as [16]

where Imetai and I 0Xide are the intensities o f the metal and oxide photoelectron peaks, N 0 and Nm are the densities of metal atoms in the oxide and metal (atoms/cm3), ^ and are theEALs of the metal photoelectron through the metal and the oxide respectively, and 6 is the electron take-off angle, defined by the surface plane to the detector.

The case for the LiF coated metals is slightly more complicated, since this uniform

overlayer model may no longer hold. However, a first approximation of the thickness can be

made by ignoring the LiF layer and using equation 6-1 directly. Subsequently, in order to

estimate the extent o f LiF coverage on the surface, the island overlayer model [22] (equation

3-12) can be used taking a ratio of the intensity of the F Is peak to the metal peak.

where duF is the thickness o f the LiF islands, and % is the area fraction o f LiF coverage.

As the oxide film grows, two configurations for the oxide structure are possible:

columnar oxides between the LiF islands or embedded LiF islands in an oxide matrix. To

determine which microstructure is likely at various stages of growth, the experimental ratio

metal(6-1)

*expm etal

J

(6-2)



of the intensity of the oxide and metal peaks can be compared with uniform and island

overlayer models modified to form the columnar and embedded models.

Columnar model:

oxide N X ee°{ l-* )ex p

T — N 2Me1 m eta l 1 V M 71 A / e X exP

- d .Q M eO

fe

\X T sine

L iF

Sind;+ ( l-^ )e x p

n M eO „ • nA Me s m v / / v Me j j

(6-3)

where dc is the thickness o f the oxide layer in between the LiF islands,

and

Embedded model:

T _ A T 2Me01 oxide “ i V o ' 1 M e 1 - exp

W

~ d cIM e O „ • r \yXMe sin 0

\ \- + ( l - ^ ) e x p) )

\ \

- dLiF '

yx™ S in 0

+ (l-^ )ex p_ ,

a L iF

(6-4)

rme,al=NMTe exp \ I X eXP .J™ Q + (l~X)^P( K v s m e | sm e j y x ™ sm u j jwhere dc is the thickness o f the oxide layer above the LiF islands, duF in this case represents both the thickness of the LiF islands themselves, and the thickness of the oxide in between the islands.

Unlike the simplification used for thicknesses with layer-by-layer oxide growth,

neither the surface coverage nor the configuration can be determined analytically using ratios

(equations 6-2 to 6-4); however, a sum of residual squares analysis can be done to fit the

angle resolved data graphically, with the minimal chi squared value indicating the best fit of

the data. The applicability o f these models is somewhat limited by the amount o f information

inherent in ARXPS spectra [23], as described in chapter 3. Therefore, some assumptions

have to be made in order to determine which model is the appropriate description o f the

oxide structure at any given time. In these extended models, the only adjustable parameter is

the oxide film thickness, with the LiF island thickness assumed from the nominal deposition

amounts, and the coverage predicted by equation 6-2 fixed for all exposures.



The thickness of the surface contamination layer was not included explicitly in this

analysis, as the attenuation should be similar for both the substrate and the overlayer;

however, an empirical constant was added to ensure adequate fitting, representing the

attenuation of the signal through the contamination layer.

In all cases, the EALs were calculated using the National Institute of Standards and

Technology EAL Database [24], which determines the EAL using the Tanuma-Powell-Penn

(TPP-2M) theoretically derived inelastic mean free path [25] and applying the transport

approximation [26], Table 6-1 lists the parameter values required to calculate the oxide

thickness, LiF coverage, and oxide film configuration.

Table 6-lParameters used for film structure analysis

M e^ M e

q M eO^ M e

2 LiF 0 L iFM e

N m NLiFF \s

K K

A1 24.01A 29.03A 33.7A 47.96A 1.5 1

6.3 Oxidation and surface structure of A1 surfaces

6.3.1 Oxidation products

Figure 6-1 below compares the evolution o f the A1 2p core level over time for uncoated and

10A LiF coated A1 surfaces. The core level shows two components, a metallic peak around

71 eV, and an oxide peak, shifted to higher binding energies by about 3eV. This oxide peak

can be attributed to hydrated AI2O3 in both cases [11]. The apparent shift in the binding

energy for the metallic component between the two systems can be attributed to differences

in surface charging, as all the spectra were aligned to adventitious C on the surface and there

is no relative shift in the oxide peak for all exposures. As the overlayer thickness increases,

surface charging o f the insulating oxide relative to the metallic underlayer becomes more



pronounced [27], increasing the apparent chemical shift. Since LiF is also an insulator

contributing to charging, there is a greater shift observed in the oxide overlayer, causing the

metallic peak to appear shifted to lower binding energies. This artefact is confirmed by the

identical O Is peak (not shown) for the coated and uncoated surfaces, implying that the same

oxide is formed in all samples.

w’c

xiS_

' ( f )cCD -*—>_cTDCDN03EoZ

Al 2p 0 .06 eV

0A L iF 1 0A L iFa ir e x p o s u re t im e

” 25 m in s

-Al 2p0.47 eV

_ a ir e x p o s u re t im e 1 5 0 0 h rs

Figure 6-1A1 2p core levels for uncoated Al and 10A LiF coated Al for exposure times of (a) 25 mins and (b) 1500 hrs. Due to the insulating nature of LiF, the coated surface shows an increasing surface charging effect with time of 0.06 eV and 0.47 eV for (a) and (b) respectively.

80 78 76 74 72 70 68 66

B i n d i n g E n e r g y ( e V )

Over time, increasing exposure to an oxidising environment causes the high binding

energy oxide component to grow for both the coated and uncoated surfaces. At room

temperature under ambient conditions, Al is expected to form a passivating oxide relatively

quickly. Uncoated Al does indeed appear to quickly reach a passivating thickness o f around

22.5A, in good agreement with other reports [11,28]. The LiF coated surface, however,

shows significantly less oxide than the uncoated metal, for all exposures, with oxide growth

continuing even after 1500 hrs. Though the 5A LiF coated samples follow the same general



trend, the amount o f oxide appears to be only slightly less than that observed on the uncoated

surface. For both 5 and 10 A coatings, the single F I s core level (not shown) was identical

and did not change with exposure, indicating the formation of stable ionic LiF on the metal

surface. The oxide thickness with exposure was determined from the ratio o f the intensity o f

the metallic component to that o f the oxide component of the Al 2p core level. For all

exposures, the fitting parameters were fixed, though slightly different for coated and

uncoated surfaces as outlined in table 6-2.

Table 6-2 XPS parameters for Al 2p core level as observed on coated and uncoated surfaces.

uncoated LiF coated

Component BE (eV) FWHM „* (eV) P

FWHM(eV) P*

Al 2p Metallic Al 71 eV§ 0.62* 3.0 0.65 3.9Hydrated A120 3 +3eV 1.58 0 1.72 0

A *f from [29] asymmetry parameter as defined in Appendix C (by p = t s * e7"-) 5After correction for LiF overlayer induced charging

6.3.2 Surface oxidation kinetics

As seen from figure 6-2 below, oxide growth appears to follow a semi-logarithmic trend as

expected by the Cabrera-Mott theory [28], with different rates for coated and uncoated Al

oxidation. Uncoated Al shows a knee in the oxidation curve, indicating the onset o f

passivation at around 60 hrs exposure. This rate is considerably slower than those reported in

some of the literature [11 30]; however, this could be attributed to the use of high purity Al

and fast evaporation rates during the deposition process. Such conditions are known to

reduce aggregation and form smooth films, which can greatly increase the time to onset of

passivation [18]. Both Chen e t a l . [11] and Hass [30] utilized mechanically polished

polycrystalline samples where fast diffusion channels for metal and oxygen ions at grain



boundaries would promote faster oxidation. The observed slower oxidation rate, however, is

a better representation of expected film behaviour for devices, as deposition conditions

similar to those used in this investigation are generally maintained during device fabrication.

Figure 6-2 Growth of oxide on Al surfaces, monitored by XPS, for thickness as estimated by the simple overlayer model. Lines represent a linear sum of reduced squares best fit of the data for the uncoated and 10A LiF coated substrates. Uncoated and 5A LiF coated Al both show a bend in the curve at around 60 hrs. The open triangles represent the predicted oxide values scaled by the LiF coverage as predicted by ARXPS.

0 1 10 100 1000

time (hrs)

With the introduction of a LiF layer, the oxidation kinetics o f the surface is changed.

With 10A LiF, there is a much lower predicted oxide thickness at all exposures, and the

oxidation rate, indicated by the change in the oxide thickness over time, is somewhat

reduced. This suggests that LiF acts as a passivating barrier on the surface of Al, significantly

modifying the oxidation rate of the metal surfaces. The thinner LiF layer has features that are

common to oxidation on coated and uncoated surfaces. Before the onset of passivation as

predicted by the oxidation o f uncoated Al, the oxidation rate is very similar. Though the

predicted thickness of the oxide appears to be lower, this is likely just an artefact due to LiF

covering some of the metal surface. Scaling the thickness values by the LiF coverage, as

described in the next section, matches up the predicted oxide thicknesses fairly well below

the onset of passivation for uncoated Al, as shown by the open triangles in figure 6-2. After

around 60 hrs, however, the scaled thickness is greater than the passivation thickness and

0A LiF26

10A LiF24

b=2.83 b = 1 .6320

b=2 .83

b = 1 .63



appears to continue to increase with time. This could indicate that the oxidation rate has

changed and is now similar to that observed on the surface of the 10A LiF layer. Although

there are very few data points for longer exposure times at 5A thickness, the behaviour can

be described, within error, by that observed on the 10A LiF coated surfaces. This suggests

that coating the Al surface with LiF does change the oxidation kinetics, but that below a

minimum amount o f LiF coverage, the impact on the overall oxidation rate of the surface is

minimal. As the thickness of the LiF layer increases, the smaller predicted oxide thickness is

no longer solely an artefact of the coverage .

Generally, the physical mechanism for logarithmic growth of oxide films is assumed

to be related to the strong electric field developed in the oxide film due to a potential

difference between metal and absorbed oxygen [28], For Al, which forms metal-excess

oxides [28,31], the oxidation rate during room temperature oxidation is proportional to the

number o f metal ions per unit area available to dissolve into the oxide, N ' , and the rate at

which these ions can escape the metal [28]. Oxidation of Al occurs primarily by outdiffusion

of the metal ions through the forming oxide layer, to the surface where they react with O2'

ions. To maintain charge neutrality within the oxide layer, the current o f positive ions is

compensated for by the transfer o f electrons to the surface to form the O2' ions.

According to the model by Cabrera-Mott and Fromhold-Cook, at low temperatures,

diffusion is driven by two mechanisms: the concentration gradient o f the metal over the oxide

layer, and the tunnelling of electrons from the metal surface. During initial oxidation, an

electric field is developed across the oxide layer due to the formation of the anions at the

2 Note that the predicted oxide thicknesses could also be attributed to 25% coverage of LiF islands on the metal surface. However, ARXPS for the nominal 10A LiF layer indicates that the real LiF island thickness would have to be less than 4A to obtain such a coverage and requires a correction factor to fit the observed results. As the relative FIs intensity for the 10A LiF layer is much greater than that of the 5 A layer, it is unlikely that the island thickness is the same with such a small increase in coverage.



oxide surface from fast tunnelling electrons. This induced field increases the diffusion of

metal atoms, allowing thin oxide scales to form quickly on the surface. Beyond a critical

thickness, the tunnelling current breaks down and electrons must be supplied by thermionic

emission. At room temperature, due to the scarcity o f thermally excited electrons, metal

diffusion is strictly controlled by the concentration gradient once the critical thickness has

been reached, and the oxidation rate is considerably slowed. Generally, if oxide growth is

logarithmic, the oxidation process can be attributed to the combination o f these two

mechanisms.

For Al, the potential barrier between the metal and the oxide, analogous to a Schottky

type barrier, is on the order o f IV [28]. This is thought to be large enough such that room

temperature is a sufficiently low enough temperature for logarithmic growth to be observed,

and the passivating oxide scale forms quickly on the metal surface. In the LiF coated system,

the oxidation would be greatly diminished as both electrons and ions are prevented from

migrating. The metal ions would have to diffuse through both the LiF lattice and the oxide

layer to reach the oxide/gas interface where oxidation takes place. Since LiF provides good

bulk lattice matching with Al over a broad range o f orientations {auf.~4.QlK [32], ciai=4. 04A

[33]), it is likely that the film will deposit along a well matched plane. The metal ions would

then have to diffuse through the commensurate LiF lattice by a self-diffusion type

mechanism without any fast diffusion channels. Due to the charge imbalance, and the

change in the size of the interstitials, the ions would have more difficulty moving through the

overlayer lattice than through the metal matrix, and the diffusivity and ion mobility

properties should be lowered. In addition, the electron injection barrier could be increased

due to the lowering o f the metal work function with the presence of LiF [34], and the

physical barrier of the LiF overlayer. With the rate o f ion diffusion, the number of sites for


1/th

ickn

ess

(1/A

)


metal to dissolve into the oxide, N ' , and the electron tunnelling probability decreased by the

overlayer, the rate of oxidation would be expected to decrease significantly.

Assuming that the Mott-Cabrera/Fromhold-Cook model holds, the oxidation rate for

the coated and uncoated surfaces can be determined from the Mott-Cabrera rate equation [28]:

r C 'A— = 2 K sinh dt

(6-5)

where x is the oxide thickness, Kox is the oxidation rate, and Cd is a characteristic distance depending on the potential electron injection barrier at the interface.

Using Ghez’s approximate solution for the rate equation [35] for oxide thicknesses

much less than the characteristic distance, x « C d , the oxidation can be described by

CdX

-In\ x 2 J

M c dK „ ) (6-6)

Figure 6-3 shows the linearized plot for the coated and uncoated surfaces. From the slope and

intercept, the oxidation rate and characteristic distances can be determined, as listed in table 6-3.

+ 0A LiFx 10A LiF0.085

0.080

0.075

0.070

0.065

0.060

0.055

0.050

0.045

7 -6 ■3 •2 1 0 1 2■5 -4

Table 6-3 Oxidation rates and characteristic lengths as determined by Mott-Cabrera theory (figure 6-3).

Kox (A/s) Q (A)

Al 2.6x1 O'9 256(±20)

LiF coated Al 2.7xl0'12 276(±18)

In(time/thickness2) (hrs/A2)

Figure 6-3 Mott-Cabrera oxidation behaviour for uncoated and 10A LiF coated Al surfaces. The solid lines represent a linear sum of reduced squares best lit of the data substrates.



According to the theory o f Mott and Cabrera [28], Q is a function of the effective charge of

a defect, eeff, the jump distance, ajump, and the contact potential, V contac t, as

no difference in the characteristic distance of the overlayer due to the presence o f LiF. As

this characteristic distance is related to the defect aiding diffusion, this suggests that the same

defect structures exist within the oxide overlayer in both cases. It also suggests that the

contact potential is the same for the two cases, and there is no change in the rate o f electron

injection. The oxidation, therefore, is predominantly controlled by the jump rate o f the ions

through the overlayer. Ion diffusion appears to be three orders o f magnitude faster in the

oxide alone compared to the combination of LiF and oxide on the metal surface. The

diffusivity of Al in AI2O3 at room temperature, extrapolated from high temperature data [36],

has a value of ~4xl0"19 cm2/s, which lies between the values for the oxidation rate o f Al for

coated and uncoated surfaces. This suggests that the oxidation may be diffusion driven,

though tracer diffusion experiments for Al ions through LiF and Al oxides formed during

exposure to the ambient environment would be of great benefit in confirming this suppressed

diffusion mechanism with LiF.

The predicted oxidation rate from the Mott-Cabrera solution is incredibly slow, much

slower than the observed oxidation kinetics on these surfaces. However, as can be seen from

figure 6-2, the calculated rate dependence requires an initial oxide thickness. As Al is known

to undergo multiple oxidation stages [37], it is likely that the behaviour being modelled here

represents a latter stage o f oxidation, just prior to the onset of passivation, where the

oxidation behaviour is controlled by ionic diffusion.

e ff jum p contact (6-7)

The similarity o f the values for the coated and uncoated metal suggests that there is



6.3.3 Surface oxide structure

With deposition of a thick LiF layer, on the order o f hundreds o f A, the overlayer could be

expected to be complete and block oxidation across the entire Al surface. With much thinner

films, angle resolved XPS analysis using the overlayer model (equation 6-2) indicates that

LiF forms islands on the metal film surface, as shown in figure 6-4.

■ data - - • 0 .9 9 , 6A — 1, 5A■ ■ - 0.05.5A m idlayer model• — 0.5, 5A• - 0.15.5A

4.5

4.02.6

3.5

3 2.5 °3 2.2

" 2.0 TUV ■

0.5

800 40 60 80 20 40 6020

take-off angle e (deg) take-off angle 6 (deg)

Figure 6-4 Determination of LiF coverage using the simple patchy overlayer model (equation 6-2). The various lines represent different values of the coverage and LiF thickness, which were the only variables used to fit the data. The close up section on the right hand side shows the predicted angular dependence with a 5A LiF layer at different coverages. The Levenberg-Marquardt reduced chi squared fit of the experimental data (dotted line) indicates coverage of 15%.

On the areas of the Al surface covered by LiF, the islands prevent oxidation, similar

to what was observed for a complete overlayer [18]. This blocking effect can be attributed to

the presumed commensurate growth o f LiF on the Al surface. Consequently, during

oxidation, the surface would have two regions with bare substrate areas of high metal ion

concentration, and LiF capped areas o f low metal ion concentration. For the diffusing

species, the LiF islands, due to lattice matching, would appear as a continuation of the Al

lattice, making the surface analogous to a corrugated metal surface.

Initially, with air exposure, the bare substrate areas become oxidized, giving rise to a

columnar structure for the film. Angle resolved XPS agrees with this columnar growth mode,



from Figure 6-5(a), with best fit for the columnar model (dotted line). The fact that the

predicted oxide thicknesses for the 5A LiF overlayer are just a factor of the coverage during

the early stages o f oxidation also supports this initial structure. If the thickness values are

scaled by a coverage o f 15%, as approximately predicted by ARXPS, they match that o f the

uncoated surface fairly well, before the onset of passivation. Over time, however, the angle

resolved data deviates from the columnar model (Figure 6-5(b)), and the apparent thickness

continues to increase.

Figure 6-5 Structure model comparisons for LiF coated Al surfaces for exposure times of (a) 25 mins with 5A LiF coverage and (b) 1500hrs exposure with 10A LiF coverage. Lines represent Levenberg-Marquardt reduced chi squared fit of the experimental data for various structure models. The solid line represents an embedded structure, the dashed line a columnar structure, and the dotted line a multilayer structure assuming a complete LiF layer at the interface.

20 30 40 50 60 70 80 90

take-off angle 6 (deg)

One possible explanation could be that once the metal ions diffuse through the LiF

lattice, oxide growth continues on the surface of the LiF islands. With such a growth pattern,

at lower surface coverages, such as those for 5A, the oxidation rate is dominated by the

oxidation o f the large metal surfaces between the islands. The observed knee in the growth

5A LiF5, 25 m'n exposure• experimental data embedded structure

— columnar structure— uniform multilayer

\ 10A LiF, 1500 hr exposure



curve (figure 6-2) could indicate the shift from oxide growth on the bare substrate to the

slower growth o f the oxide layer on top o f the islands. This is also supported by the

overestimation and apparent continued growth of the oxide thickness after passivation.

At higher coverages, at least 60% as observed for the thicker LiF film, the oxidation

process is dominated by Al ion diffusion limited growth through the LiF layer from the

earliest stages. In this case, the apparent decrease in the thickness is not an artefact o f the

coverage.

Schematically, oxide growth on LiF capped Al can be described as a multistage

process, resulting in an embedded structure o f the oxide film, as in Figure 6-6.

(a)

m .

oxide

V//A

Al substrate

(b)

oxide

LiF LiF LiF

Al substrate

Figure 6-6 Schematic oxide growth model on LiF coated Al surfaces, (a) Initially, growth occurs between LiF islands, producing a columnar structure. As growth progresses, Al diffuses through the LiF islands and growth occurs over the islands, leading to (b) an embedded structure.

6.3.4 Impact o f LiF on metal surface oxidation

The final thickness of the oxide layer is set by the potential field built up at the

surface of the oxide. Above the critical thickness, ion migration is no longer accelerated by

an electric field, and oxidation is basically halted [28]. If the embedded model o f the oxide

structure can be assumed, and the electron tunnelling is similar for the two cases, the total

physical thickness of oxide formed for Al should be the same regardless of the surface



condition. Given that the difference in the oxidation rate is three orders o f magnitude, one

could predict that LiF coated surfaces would reach a fully oxidized thickness after 104 hours,

compared to 60hrs for uncoated Al surfaces. This implies that LiF can delay the oxidation

both by reducing the surface area and by changing the oxidation kinetics. Substantial

lifetime improvements have in fact been observed by others with extremely thick LiF layers,

250A and thicker, which showed no visible oxide growth even after months o f exposure in

ambient conditions [19].

6.4 Estimation of device failure due to oxidation of Al/LiF based cathodes

In a device structure, the oxygen load is considerably lower than that for metal

surfaces exposed to the ambient environment. Therefore, the impact o f the change in the

oxidation kinetics should be more pronounced. Since cathode oxidation is a major failure

mechanism in OLEDs [38,39], device shelf time should be influenced by the interlayer

thickness. With minimal stress on the devices, the mean time to failure due to interfacial

oxidation can be determined. Devices with a C6o electron transport layer are best suited for

this type of analysis, as it has been reported that the A1/C60 contact will degrade from an

ohmic contact to a blocking one after exposure to air due to the emergence of a potential

barrier between the top electrode Al and C60 film [40]. In addition, the interfacial structure

can also be examined by peel-off to see the extent of oxidation, without interference from the

Al3+ component of Alq3 in the Al 2p core level.

The interlayer thickness did have an effect on the shelf time of the C60 based LiF/Al

bilayer cathode devices. Initially, as expected, the LiF thickness modifies the device

performance slightly, with the 5A LiF interlayer showing the best device properties. In order



to check the device performance as a function of shelf time, the devices were stressed in air

only slightly, to 5V, to measure the relative change in the properties o f each device over

time. The minimal stress from running the experiment had no effect on the observed device

properties, with similar device performance for stressed and unstressed devices after 7 days

(figure 6-7). The only exception was the 20A LiF device, which no longer shows appropriate

diode behaviour after the 2nd day. It is likely that this was just a bad device which shorted out

after the voltage was applied. This is confirmed by the behaviour o f an unstressed device at

that thickness on the seventh day, which showed good diode behaviour and a similar

decrease in performance after the seventh day as the other thicknesses without failure. For

the device with 20A LiF, therefore, the behaviour over time was determined from the original

data for the first and second day and the data for the seventh day from the unstressed device,

3.0

10A LiF5A LiF

first day last day (stressed) last day (unstressed)

20A LiF 30A LiF

V oltage

1 2 3 4 5 6

V oltageFigure 6-7 Comparison of device behaviour on the first and 7th (final) day of the experiment. The circles and triangles indicate the behaviour of stressed devices and unstressed devices after the same length of exposure, indicating similar behaviour. The behaviour of the stressed device with 20A LiF no longer shows appropriate diode behaviour after the 2nd day of stressing, but the unstressed device on the final day indicates a similar trend as for all the other thicknesses.



90'(b)

0 1 2 3 4 5 6 7 8 9

Shelf time (days)5 10 15 20 25 30

Nominal thickness (A)Figure 6-8 (a) Current decay measurement for C60 based devices with Al/LiF cathodes of varying LiF thickness. In region (1), the performance decays in relation to the LiF thickness. After one day of exposure, region (2), the device performance decays exponentially, with the same decay constant. The solid lines are guides to the eye, but in region (2) indicate the reduced squares best fit of exponential decay with a decay constant as determined from figure 6-9. The device with a 20A LiF layer has very similar exponential decay behaviour to that of the 30A LiF device, (b) The thickness dependent percentage decrease in current after the first day.

Figure 6-8(a) below shows the decay of the maximum achievable current as a

function of the shelf time. In all cases, the current decreases rapidly after the first day, then

appears to decay exponentially. This behaviour can be explained by the oxidation behaviour

of the cathode. Initially, the rapid decrease in the maximum achievable current is dependent

on the nominal deposited thickness, which presumably changes the protection coverage of

the cathode, i.e. the LiF covers more of the organic surface at thicker depositions, protecting

the cathode deposited on top from oxidation. The thinnest interlayer, 5A LiF, provides the

least protection against cathode oxidation, and the device performance decays by nearly 90%

almost immediately. As the LiF layer thickness increases, the percentage decay in the current

decreases, as shown in figure 6-8(b).

After this initial decrease, the device properties appear to follow an exponential trend.

The decay constant, xa, can be determined from



J J,— = —-exp

J.

t

' o u o V ^ d J

where J0 is the initial measured current density, .7/ is the measured current density after 1 day, and t is time.

As figure 6-9 shows, the data can be well described by a single value for the decay

constant, suggesting that the mechanism behind device degradation is the same for all

devices. As this exponential decay only begins after the first day, to determine the decay

constant, the maximum current was renormalized with respect to the ratio after one day of

exposure. The device with a 20A interlayer was not completely consistent with these results,

since it could be considered to have failed after the 2nd day; therefore, the decay constant was

determined without that data.

(6-8)

Figure 6-9 Renormalized maximum current achievable over time. The solid line represents a linear sum of reduced squares best fit of the data.

0 1 2 3 4 5 6 7 8 9

Shelf time (days)

We could tentatively assign exponential decay behaviour to the cathode oxidation

with LiF presence, even though the oxidation rate predicted in section 6.3.4 is much slower

than this decay constant. It is impossible to unambiguously assign this decay constant to Al

oxidation with LiF protection, as the decay rate is an order o f magnitude greater even than

the empirical values derived from the oxidation kinetics of LiF coated Al surfaces.

♦ 10A ► 30A x average

0.0

-0.5

0

5



Conversely, complete oxidation of the cathode surface could all be occurring within the first

day, the amount o f which is affected by the thickness o f the LiF layer. In such a case, the

exponential decay of the device behaviour after the first day, which is the same for all

devices, would no longer be related to cathode oxidation. Nonetheless, the implication is that

LiF blocks regions of the surface from oxidation, preserving device properties, with the

amount of blockage related to the thickness o f the LiF layer.

If device failure is taken as a 90% decrease in the current (since these devices are not

encapsulated), then the shelf time can be described as a function of thickness, combining

both the effects of interlayer thickness and the exponential decay due to oxidation, as in

figure 6-10. Since the decay in device properties is related to the cathode oxidation, this shelf

time dependence is mostly a reflection of the protection provided by the LiF layer. A 5A LiF

provides very poor protection for the Al cathode, as was also observed for coated Al

surfaces, and the device does not last more than one day. A slightly thicker layer improves

the shelf time by nearly a factor of three, reflective of the noticeable passivation o f the Al

surface with 10A of LiF. Above 20A, the shelf time becomes independent o f the LiF

thickness, suggesting that complete coverage at the Al interface with LiF has occurred.

8 ■

Figure 6-10 The shelf time of the devices, delined as the maximum time to reach 10% of initial device performance. The lines are just a guide to the eye.

a>E

<DJZCO

10 15 20 25 305

Nominal thickness (A)



If these devices had been encapsulated, as is the current practice for commercial

OLED products, the impact o f the LiF in preventing cathode oxidation would be greatly

enhanced. Current requirements for commercial OLED encapsulation limit the amount of

moisture penetration to 10pg/m2/day [38, 39]. As the devices were stored and tested without

any encapsulation, this investigation represents an accelerated test o f device aging, with

exposure to 4.84x104 jig/day of moisture. For a device with 1mm2 area, assuming that 20%

of the moisture outside the device can reach the cathode/organic interface by diffusion

through the grain boundaries in the cathode, and laterally through the device, each day of

shelf time for the exposed device is representative o f nearly 2 % years for an encapsulated

device. Therefore, a device with more than 20A LiF would last nearly 18 years on the shelf

before degrading to 10% of its initial performance when manufactured.

6.5 Interfacial chemical structure at the Al/LiF/organic interface

There appears to be some reduction o f cathode oxidation as the thickness of the LiF

layer increases, but it is not totally suppressed, consistent with the oxidation behaviour of

coated Al surfaces. Figure 6-11 (a) shows the Al 2p core level at the cathode surface after

peel-off o f the cathode/organic interface for both Al and A1/100A LIF cathodes with C6o as

the electron transport layer. In both cases, two components are visible, likely due to a

metallic and an oxidized component o f the core level. For the pure metal cathode, there is

more visible oxidation at the organic interface, as shown by large high binding energy core

level (figure 6-11(a)). With the LiF interlayer, the Al 2p core level is less visible, but appears

to have significant oxidation. The amount o f oxygen (not shown) is similar for the two cases.

As the bi-layer system with C6o can be considered as a Metal-Inorganic-Metal (MIM) type


Inte

nsity

(a

rb.

units

) In

tens

ity

(arb

. un

its)


electrode3, the metal/LiF surface is another essential contact, as it regulates the amount of

charge carriers that are injected into the LiF layer, where they can be stored. The device

properties are, therefore, likely also controlled by the oxidation at the metal surface. Figure

6-11(c) indicates that at the metal surface, there is greater oxidation, and three times the

amount of oxygen, at the metal surface without a LiF interlayer.

Al 2p i

‘A

I ,

A l 2p JII

(C )-- • - A l f' — Li F/ AI J

I* "

. j. . .f. -

• . . /* * ? 'v i?V# ff 7 l * ■ L\ •1 «i . i . i ,

1W i78 76 74 72 70

Binding Energy (eV)6 8 80

(d)metallic cathoiAl oxide

78 76 74 72 70Binding Energy (eV)

78 76 74 72 70 68

Binding Energy (eV)

Figure 6-11 Al 2p core level for (a) Al surface (b) A1/100A LiF surface after peel-off at the cathode/organic interface (c) the metal surface of the cathode after Ar+ sputtering (d) the sputter profile through the thickness of the LiF layer for Al/IOOA LiF cathodes showing the evolution of the chemisorbed Al.

In a device using a multilayer cathode with LiF, the oxidation suppression effect of

the LiF is magnified, as described above, since there would no longer an abundant supply of

oxygen for oxidation, as was available for metal surfaces. For the device with a C6o layer

sandwiched between electrodes 1mm across, the oxidation would be strictly limited by lateral

- 13 2diffusion of O through the organic layer. With a diffusion constant on the order o f 10“ cm /s

for C6o [41], there would only be 0.3% of the O at the metal surface 120 pm laterally inside

the device after 1500 hrs. Even with grain boundary and pinhole diffusion through the

See Chapter 8.



cathode thickness, the oxygen load at the buried interface would be expected to be low. With

such little oxygen, even small amounts o f LiF could effectively prevent complete oxidation

o f the cathode. At much larger thicknesses, though the LiF completely covers the organic

surface, the cathode is not completely free o f oxidation. The lattice coherence o f the

interlayer and the cathode does not prevent some intermixing of the layers. Metallic and

oxidized Al visible in XPS throughout the LiF thickness, suggests that Al might be diffusing

through the LiF layer.

The change in the relative amount of Al with chemisorbed O through the thickness of

the cathode with the LiF interlayer is shown in figure 6-11(d). During deposition, as the Al

ions diffuse through the interlayer, they encounter laterally diffusing oxygen atoms, and trap

them away from the injection zone. This has the additional benefit o f consuming some

oxygen that could act as bulk conduction traps within the C6o layer itself [20], Therefore, the

LiF interlayer encourages conduction by both scavenging oxygen within the LiF layer and

preventing oxidation at a critical injection region.

When devices are allowed to degrade with exposure to air, the diode performance of

devices with even a thin LiF interlayer can be recovered with annealing [20], As figure 6-12

shows, a device with a LiF interlayer showed similar injection in forward and reverse bias.

Without a LiF layer, the device had very little forward injection. The interlayer, therefore,

preserves the device characteristics by blocking Al oxidation at the metal surface. As

discussed by Huang et al., annealing of the devices liberates the trapped O within the organic

and LiF layers and the device shows similar forward and reverse bias performance. Without

the LiF layer, most of the oxygen is trapped at the metal surface as oxides, and cannot be

removed by annealing.



Figure 6-12I —V characteristics of the C6o sandwich diodes with Al and Al/LiF electrodes after exposure to air for 1 hr and then baked in vacuum for 24 hrs (excerpted withpermission from [20],Copyright 2005, American Institute of Physics).

- 2 - 1 0 1 2

Voltage (V)

The organic layer appears to have little impact on the oxidation behaviour, with

cathode oxidation through the thickness showing a similar trend for the A1/100A LiF/Alq3

combination. It is slightly more difficult to judge the oxidation behaviour o f Al/Alq3 systems

as the oxidation state of Al is very similar for the oxide and the molecule. However, there is

less visible oxidation at the interface, once the molecular layer has been removed, as Alq3

tends to react with oxygen and moisture, and crystallize [42], O diffusion through the device

may have been prevented by reactions with the molecules at the edges of the device.

with LiF without LiF

0.003

0.000

-0.003

A l 2 p O l s F 1s

ALO. 32nm

d=0nm

80 76 72 68 404 400 396 536 532 528 692 688 684 680Binding Energy (eV)

Figure 6-13 Depth profile results for a 200A LiF layer, showing the complete blocking of oxygen diffusion from the organic layer. The residual Alq3 is removed after the first two cycles. AIq3 shows very little lateral diffusion of oxygen, so oxidation of metal surface due to diffusion from outer cathode surface through grain boundaries or during initial deposition[43[.



With 200A LiF, the oxidation is almost completely blocked, as shown above in figure

6-13, and described in Chapter 9. A small amount of oxygen visible within the LiF layer

itself, can be attributed to diffusion through the organic or the LiF layer itself. The Al 2p core

level again shows small amounts o f chemisorbed O at the interface with LiF; however, most

o f the oxidation o f the Al cathode appears to have occurred mainly during deposition or from

oxidation of the top o f the cathode rather than at the buried interface, as the amount o f oxide

does not change substantially beyond the LiF layer.

Therefore it is likely that above 20A LiF, where the coverage o f the organic layer is

complete, the cathode oxidation is controlled by oxygen and water vapour diffusion through

defects in the cathode, as well as lateral diffusion through the organic and LiF layers. The

LiF layer, therefore, prevents oxidation at a critical injection surface and protects the device.

6.6 Summary

At nominal thicknesses typically used in optoelectronic device cathodes, deposited

LiF does not completely cover the surface, but forms islands. For Al, even without complete

surface coverage, LiF is effective in slowing down oxidation due to the lattice matching of

the overlayer and the substrate as predicted from bulk lattice constants. Hence, the

commensurate LiF islands give the Al surface a corrugated structure, upon which the oxide

grows, with the islands acting as diffusion barriers for Al atoms. 10 A LiF (61% coverage) is

sufficient to significantly modify the oxidation kinetics, due to an ion diffusion dominated

oxidation mechanism.

These changes in the oxidation of the coated surface can explain the degradation of

organic light emitting devices with reference to the interfacial oxidation. The observed



suppression of oxidation of a multilayer structure in air would be magnified within a device,

where the oxygen load is greatly reduced. With a lattice matching interlayer, such as LiF

with Al, the improved oxidation resistance at the interface may explain the increased shelf

times observed with the use o f multilayered cathodes.

6.7 References

1 C. W. Tang and S. A. Van Slyke, Appl. Phys. Lett. 51, 913 (1987).

2 E.I Haskal, A. Curioni, P.F. Seidler, and W. Anderioni, Appl. Phys. Lett. 71, 1151 (1997).

3 V-E. Choong, M. G. Mason, C. W. Tang, and Y. Gao, Appl. Phys. Lett. 72, 2689 (1998).

4 N. Johansson, T. Osada, S. Stafstrom, W. R. Salaneck, V. Parente, D. A. dos Santos, X.Crispin, and J. L. Bre'das, J. Chem. Phys. I l l , 2157 (1999).

5 T. Osada, P. Barta, N. Johansson, Th. Kugler, P. Broms, and W.R. SalaneckSynth. Met.102, 1103 (1999).

6 A. Turak, D. Grozea, X.D. Feng, Z.H. Lu, H. Aziz, A.-M. Hor, Appl. Phys. Lett. 81, 766 (2002).

7 V. Choong, Y. Park, Y. Gao, T. Wehrmeister, K. Mullen, B. R. Hsieh, and C. W. Tang,Appl. Phys. Lett. 69, 1492 (1996).

8 L T. Wakimoto, Y. Fukuda, K. Nagayama, A. Yokoi, H. Nakada, and M. Tsuchida, IEEE Trans. Electron Devices 44, 1245 (1997).

9 M. G. Mason, C. W. Tang, L-S. Hung, P. Raychaudhuri, J. Madathil, D. J. Giesen, L. Yan,Q. T. Le, Y. Gao, S-T. Lee, L. S. Liao, L. F. Cheng, W. R. Salaneck, D. A. dos Santos, and J. L. Bredas, J. Appl. Phys. 89, 2756 (2001).

10 C. Chen, S. J. Splinter, T. Do, and N. S. McIntyre, Surf. Sci. 382, L652 (1997).

11 N. S. McIntyre and C. Chen, Corr. Sci. 40, 1697 (1998).

12 V. Fournier, P. Marcus, and I. Olefjord, Surf. Interface Anal. 34, 494 (2002).

13 T. Do, S. J. Splinter, C. Chen, andN. S. McIntyre, Surf. Sci. 387, 192 (1997).

14 J. van den Brand, W. G. Sloof, H. Terryn, and J. H.W. de Wit, Surf. Interface Anal.36, 81 (2004).


16 J. E. Gray and B. Luan, J. Alloys Compounds 336, 88 (2002).

17 P. H. Beming, G. Hass, and R. P. Madden, J. Opt. Soc. Amer 50, 586 (1960).

18 D. W. Angel, W. R. Hunter, R. Tousey, J. Opt. Soc. Amer 51, 913 (1961).

19 J. T. Cox, G. Hass, and J. E. Waylonis, Appl. Optics 7, 1535 (1968).



20 C. J. Huang, D. Grozea, A. Turak, and Z. H. Lu, Appl. Phys. Lett. 86, 033107 (2005).


22 T. A. Carlson, Surf. Interface Anal. 4 , 125 (1982).

23 P. J. Cumpson, J. Electron Spectrosc. Relat. Phenom. 73, 25 (1995).

24 C. J. Powell, NIST Electron Effective-Attenuation-Length Database (National Institute of Standards and Technology, Gaithersburg, MD, 2001), p.Version 1.0.

25 S. Tanuma, C. J. Powell, and D. R. Penn, Surf. Interface Anal. 21, 165 (1994).

26 A. Jablonski and S. Tougaard, J. Vac. Sci. Technol. A 8, 106 (1990).27 T. L. Barr, in Practical Surface Analysis, 2nd edition, edited by D. Briggs and M. P.

Seah (Wiley & Sons Ltd, New York, 1990), Vol. 1, Chap. 8, p.370

28 N. Cabrera and N. F. Mott, Rep. Prog. Phys. 12, 163 (1948).29 Fundamental XPS Data from Pure Elements, Pure Oxides XPS International Inc. 1999

30 G. Hass, Z. Anorg. U. Allgem. Chem. 254, 96 (1947).31 D. A. Jones, Principles and Prevention o f Corrosion, 2nd edition (Prentice-Hall, Upper

Saddle River NJ, 1996), Chap. 12, p.412-417.

32 R. N. Euwema, G. G. Wepfer, G. T. Surratt, and D. L. Wilhite, Phys. Rev. B 9, 5249 (1974).33 W. B. Pearson, A Handbook o f Lattice Spacings and Structures o f Metals and Alloys

(Pergamon Press,, New York, 1967), Vol. 2.

34 R. Schlaf, B. A. Parkinson, P. A. Lee, K. W. Nebesny, G. Jabbour, B. Kippelen,N. Peyghambarian, and N. R. Armstrong, J. Appl. Phys. 84, 6729 (1998).

35 R. Ghez, J. Chem. Phys. 58, 1838 (1973).36 P. Kofstad, Nonstoichiometry, Diffusion and Electrical Conductivity in Binary Metal

Oxides (Wiley and Sons, Inc., New York, 1972), p. 343-348.

37 S. A. Flodstrom, R. Z. Bachrach, R. S. Bauer, and S. B. M. Hagstrom, Phys. Rev. Lett.37, 1282 (1976).

38 P. E. Burrows, V. Bulovic, S. R. Forrest, L. S. Sapochak, D. M. McCarty, and M. E. Thompson, Appl. Phys. Lett. 65, 2922 (1994).

39 J. S. Lewis and M. S. Weaver, IEEE J. Quantum Electron. 10, 45 (2004).

40 H. Yonehara and C. Pac, Appl. Phys. Lett. 61, 575 (1992); C. H. Lee, G. Yu, D. Moses, A.J. Heeger, and V. I. Srdanov, Appl. Phys. Lett. 65, 664 (1994).

41 B. Pevzner, A. F. Hebard, M.S. Dressselhaus, Phys. Rev. B 55, 16439 (1997).

42 H. Aziz, Z. Popovic, S. Xie, A-M. Hor, N-X. Hu, Carl Tripp, and G. Xu, Appl. Phys.Lett. 72, 756 (1998).

43 D. Grozea, A. Turak, X.D. Feng, Z.H. Lu, D. Johnson, R. Wood. Appl. Phys. Lett. 81, 3173 (2002).


Chapter 7

LiF/metal bilayer structures II - Case of Mg/LiF

7.1 Introduction

As described in chapter 6, interfacial oxidation behaviour has a significant impact on

subsequent device performance. As many cathode metals strongly interact with the organic

layers, one of the proposed mechanisms for performance improvement with LiF interlayers

has been the suppression o f interfacial breakdown reactions between the reactive metals and

the susceptible organic active layers [1]. In order to understand the potential interfacial

reactions with such multilayer structures in the device, it is important to first understand the

oxidation behaviour of significant metal-interlayer combinations under ambient conditions.

Most of these multilayer cathodes use standard industrial materials such as Al and Mg:Ag

- 124 -


Chapter 7 LiF/metal bilayer II: Case of Mg/LiF 125

alloys as a primary cathode component. Whereas Al/LiF cathodes show improvements in

device performance and shelf time related to the prevention o f interfacial oxidation, Mg/LiF

cathodes show particularly poor device characteristics [2,3], suggesting a different interfacial

mechanism. As Mg also has a well described oxidation behaviour, LiF/Mg surfaces represent

ideal systems for a study o f multilayer cathode oxidation and the impact o f LiF.

Much research has been carried out on Mg oxidation [4,5,6,7,8,9], especially

regarding the sensitivity o f the oxidation processes to small surface activity changes.

Deliberate surface activity modification with overlayer coatings is widely utilized, for both

passivation and activation o f metal surfaces [10]. On Al, which self passivates at room

temperature, overlayers generally decrease the oxidation rate by forming a physical barrier

on the surface. Similarly, thick MgF2 layers have been used to protect Mg surfaces from

water uptake during metal electroplating[10]. Mg, however, is also prone to surface

activation by surface impurities. McIntyre and Chen [11],for example, found that Mg alloys

had much thicker oxides and degraded considerably faster than pure Mg, due primarily to

galvanic coupling between inclusions and Mg in the presence of water. With extremely thin,

5-10A, layers o f LiF as used in organic optoelectronics, the probability o f a complete

interlayer blocking the surface is unlikely, and the impact of such thin overlayers on the

oxidation kinetics has been largely unexplored.

In this chapter, to clarify the effect of LiF at Mg interfaces, we describe the use of

XPS to monitor the oxidation kinetics and by-products on Mg surfaces coated with thin

layers of LiF, and in devices using Mg/LiF cathodes. The observed activation o f Mg surfaces

suggests that lifetime improvements in devices can be tailored by controlling the cathode

surface activity with “lattice matching” interlayers.


Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF

7.2 Experimental

126

The coated metal structures for surface oxidation were produced using the Kurt J. Lesker

OLED cluster tool by thermal evaporation o f 5000A of Mg onto Si (100) substrates under

10'6 Torr vacuum. Shadow masks were then used to deposit thin layers of LiF on half the

surface at a rate of 0.5 A /s. Samples nominally coated with 5 and 10A of LiF and uncoated

metal surfaces from the same wafer substrate were then exposed ex-situ to laboratory air

(300K, 20-30% relative humidity) for various times ranging from 20 mins to 3000 hrs.

Functional OLEDs were fabricated using the procedure described in Chapter 4. In this

case, all layers were deposited in OLED cluster tool for a structure as follows:

glass/ITO/600A NPB/400A Alq3/2000A Mg. Half o f the pixels also included a 10A LiF layer

between Alq3 and the Mg cathode. Some samples were peeled-off in-situ and examined by

XPS. The performance o f devices deposited on the same glass substrate was also measured.

Spectra were generated by a monochromated A1 Ka (1486.7eV) source with a

23.35eV pass energy. The photoelectron take-off angle was varied between 25° and 85° for

the angle resolved analysis of the surface. Least-squares curve fitting analysis was carried

out as described in chapter 4. The shape and area of the metallic core level was kept constant

for all exposure times.

7.3 Oxidation products and kinetics of Mg surfaces

The oxidation kinetics o f Mg are expected to differ from those o f Al, as the oxidation of Mg

is controlled by the partial pressures and diffusion rates of atmospheric gases through the

oxide, rather than the diffusion of metal ions. In addition, over time the surface oxide will



partially convert into hydroxides and carbonates due to exposure to water and carbon dioxide

under ambient conditions [4-6]. Figure 7-1 shows the evolution of the oxidation products o f

coated and uncoated Mg surfaces. For the coated Mg surfaces, the high binding energy

shoulder grows and shifts to higher energies with increasing exposure, compared to the

uncoated surfaces. The Mg 2p core levels shown here have been aligned to metallic Mg at

49.5eV [12], because alignment with adventitious C did not line up the high energy peaks

from the oxidation products, unlike the case o f A1 described in chapter 6. Alignment to C Is

for MgO surfaces generally underestimates the binding energy o f the metallic peak due to the

interaction of C with the F-centers in MgO and is therefore an unreliable standard when

comparing oxides produced under different conditions [13], In this study o f Mg/LiF, the

metal films were deposited and oxidized under the same conditions. Therefore, the

■ M g 2p ' 0- a i r e x p o s u r e t i m e —■ 7 . 8 h r s J

i f \! / M g l M g<

._ _ i_ _ i_!_ i_ _ i_ _ i_

1 e V ( a ) :OALi F.

7 1 — 1 o AL i F -

M 9 ° | J

0 H ) 2 l J 0 \

'■ ; i 1 ' j ■ i . ' i ■ i• Mg 2p 0 .4 7- a i r e x p o s u r e t im e _ _' 1 5 0 0 h r s f

! i tJ Jm&o M g 0- . . . . . . . . . . . . . . i . i

e V “ " ' ( W

^ -W 11

M g l :

H )2 \ : . i .

Figure 7-1 Mg 2p core level for uncoated Mg and 10A LiF coated Mg for exposures times of(a) 7.8hrs and (b) 1500hrs. Both uncoated and coated surfaces show a pronounced high binding energy shoulder, corresponding to a superposition of Mg(OH)2 and MgO states. For the LiF coated surface, there is a shift of 0.1 and 0.47eV due to surface charging for (a) and (b)respectively.

56 55 54 53 52 51 50 49 48 47 46

Binding Energy (eV)



observable difference in the position and shape of the high BE shoulder for the coated and

uncoated Mg surfaces indicates a change in the surface activity o f the coating, and the

possible formation o f different compounds on the two surfaces.

The separation between the metallic and oxide peaks, and the FWHM are

substantially greater for the LiF coated Mg surfaces than uncoated surfaces. Part o f the

binding energy difference can be attributed to the same overlayer charging effect as was

observed for LiF coated A1 surfaces, but this would only account for half of the observed

difference, as shown in figure 7-2(a). For both coated and uncoated surfaces, the high

binding energy component broadens with exposure, indicating slight charging of the oxide

overlayer and increasing oxide thickness; however, the FWHM for the coated surface is

higher, and broadens more appreciably as exposure increase, as shown in figure 7-2(b).

. . . .- o - charging shift from LiF

- - M9o»de-M9°

o ’ ’ OALiF ♦ 10ALiF

2.6 (1500,2.5)-1.4

2.51.2

2.4

1.0 2.3

> 2.20T2.1> ° ' 80

LU0600< 0.4

(1500,1.97;X 2.0

LL

0.2

0.0

10 100 1000100 100010

Exposure time (hrs) Exposure time (hrs)

Figure 7-2 (a) The binding energy difference between the most intense peak from metallic Mg and that from the higher binding energy side of the Mg 2p core level. Open diamonds represent the shift in the hydroxide binding energy from charging due to the presence of LiF as deduced from the shift to the A1 oxide peaks in chapter 6. Lines are just a guide to the eye (b) The change in the FWHM of the high binding energy component of the Mg 2p core level. The lines represent a linear sum of reduced squares best fit of the data.

This increase in the FWHM, along with the increasing asymmetry of the high binding

energy component with exposure toward the low binding energy side, suggests the presence

of a second oxidation product in the case of coated surfaces, which are not observed for


Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 129

uncoated surfaces. Figure 7-3(a) shows the fit of the Mg 2p core level for 10A LiF coated Mg

with three components, one metallic state and two oxidation states, after 1500 hrs exposure.

The curve fitting was performed for all exposures keeping the fitting parameters fixed for the

metal, as shown in table 7-1. As well, it was assumed that the first oxidation state visible in

the spectrum was the same as that for the uncoated Mg film. Since the LiF coating would

induce surface charging on Mg surfaces similar to Al, the position of this first oxidation state

in the Mg 2p core level was corrected for LiF overlayer charging, as indicated in figure 7-2

above.

. Mg 2p

55 54 53 52 51 50 49 48 47co

'c13

-Q

COcCD

' 1 i i----1----*... r 1 i » ... i

' c 1s A(b).

adventitious* \C- ° H j \

A W \J \J adventiticaus/ MgCOg c -C

1 . 1 ......................... ......c

O 1sMg(OH),

MgOM^CO

538 536 534 532 530

Figure 7-3 (a) Curve fitting results for Mg 2p of 10 A LiF coated Mg at 1500 hrs exposure. The experimental data (open diamonds) can be well fitted by the sum (solid line) of three separate peaks (dashed lines), one metallic state at 49.5eV, one hydroxide/oxide state at 51eV , and a carbonate state at 52eV. Oxide values include a 0.47eV charging offset due to the insulating nature of LiF on the surface of Mg.

(b) C Is core level of 10A LiF coated Mg (open diamonds) surfaces after 1500 hrs exposure, with three chemical states attributable to adventitious C (284.6eV and 286eV) and M gC03 (289.5eV). There may also be a slight contribution at 291eV, also likely due to adventitious C.

(c) Curve fitting results for O Is of 10 A LiF coated Mg at 1500 hrs exposure. The experimental data (open diamonds) can be well fitted by the sum (solid line) of three separate peaks (dashed lines), an oxide state at 531.3eV, a hydroxide state at 532.9eV, and a carbonate state at 533.9eV. There is likely also a contribution from the adventitious C-OH beneath the carbonate peak at 531.3eV that could not be resolved.

Binding Energy (eV)



With long term exposure, Mg generally oxidizes into a mixture of Mg(OH)2 and

MgO, which show no resolvable binding energy difference in the Mg 2p core level [6];

therefore, the first high binding energy peak at around 50.8eV can be attributed to a

superposition of the MgO and Mg(OH)2 states. The second oxidation state in the Mg 2p core

level for coated substrates, with an ~1 eV chemical shift greater than MgO, can therefore be

attributed to MgC0 3 [6]. The C Is core level (Figure 7-3(b)) confirms the presence of this

carbonate structure with the shoulder visible at 289.5 eV [6, 12]. All three components are

also clearly visible in the O Is core level (Figure 7-3(c)), with a dominant hydroxide peak at

533.leV, and visible low and high binding energy shoulders for MgO at 531.1eV and

MgC0 3 at 534.3eV [5]. The binding energies for the hydroxide and carbonate peaks are

slightly higher than that observed previously, but are similar for coated and uncoated films.

By 1500 hrs exposure, the uncoated Mg films also begin to show evidence of some carbonate

formation. The binding energies and fitting criteria for the various components are

summarized in table 7-1.

Table 7-1 Summary of peak positions and curve fitting parameters for coated and uncoated surfaces of Mg

Corelevel Component BE (eV)

uncoated

FWHM (eV) P* BE(eV)

LiF coated

FWHM P*

Mg 2p Metallic Mg 49.5 0.58*4 Increases with

2.8 49.5 0.64

Increases with1.4

MgO +1.3±0.1 exposure(1.73-1.97)

0 +1.3±0.1§ exposure(1.36-1.78)

0

MgC03 - - - +2.3±0.1§ 1.50 0C Is MgC03 - - - 289.5 1.36 0O Is MgO 531.1 1.61 - 531.1§ 1.60 0

Mg(OH)2 533.1 1.73 - 533.1§ 1.58 0MgC03 - - - 534.3§ 1.95 0

i * —f from [14], asymmetry parameter as defined in Appendix C (by p = TS*eTL) s After correction for LiF overlayer induced charging



The impact o f multiple oxide and carbonate formation on the LiF coated surface is

evident as well in the growth of the oxide coating over time, as shown in figure 7-4. There is

a noticeable change in the growth curve for the LiF coated surface after lOOhrs, where the

oxide growth rate increases drastically. Therefore, the LiF coated Mg surface, unlike

uncoated Mg, shows two distinct oxidation regimes. ARXPS again shows the existence of

LiF islands on the surface, with about 65% coverage as in the A1 case. With such island

coverage, initially, the LiF layer would act as a physical barrier, blocking a portion of the

surface with less oxide formation overall. During this regime, MgCCL is the dominant

component of the coating, as shown in Figure 7-4(b). As oxidation progresses, hydroxide

formation starts to increase then become dominant as the overall amount of oxide increases

dramatically.

Figure 7-4 Growth of oxide on Mg surface monitored by XPS (a) for uncoated Mg and 10A LiF coated Mg surfaces. Lines represent a linear sum of reduced squares best fit of the data. 10A LiF coated Mg shows a bend in the curve after 100 hrs. (b) Growth of various oxide components for the LiF coated surface. The onset of the bend observed in (a) corresponds to a shift from carbonate dominated growth to hydroxide dominated growth. The dotted lines are just a guide to the eye.

10 100 1000

Time(hrs)

o 0A LiF ♦ 10A LiF

0.9

0.8

0.5

Mg(OH)2+MgO u p

▼ M gC030.2

0.0

o>- 0.2

o>-0 .4

- 0.6



Mg forms anion-vacancy oxides, with oxide growth proceeding by field induced

outward migration o f vacancies rather than o f ions, i.e. the inward diffusion of oxygen to the

oxide/metal interface [11], As well, the conversion of MgO into hydroxide and carbonate

phases is limited by the water vapour and CO2 partial pressures. With LiF covering portions

of the surface, the thermodynamically preferred carbonate phase [15] would quickly form,

with the rest of the oxide converting to hydroxide. By contrast, the uncoated Mg surface has

a greater proportion o f unconverted MgO in the same ambient environment, as in Figure 7-5.

Figure 7-5 O Is core level for 10 A coated and uncoated Mg surface after 1500 hrs exposure, indicating a greater amount of unconverted MgO for uncoated Mg. For the LiF coated surface, there is a shift of 0.47 eV due to surface charging.

540 538 536 534 532 530 528

Binding Energy (eV)

As with Al, the oxidation for LiF capped surfaces is diffusion limited with increasing

coverage, in this case requiring atmospheric anion migration through the LiF lattice.

However, since LiF and Mg are not well lattice matched over a number of crystal

o o

orientations ( a /c Mg = 3.2 A/5.2 A [16]), there would be some fast diffusion pathways that

enhance oxidation at the island edges. LiF [17] has lower adsorption energies for carbon

oxide species than MgO [18] and a high affinity for C, interacting with conjugated C species

to form charge transfer complexes [19]. On the heterogeneous surface, the edges of the LiF

islands would act as prime nucleation centers for carbonate growth, encouraging both

0.47 eVO 1sair exposure time 1500 hrs

w - OALiF— 10A LiFc

Z!

CO

■£>wc0) -I—»cTJ0N

MgO

coEoz



uncovered substrate oxidation, and uptake and trapping o f C at the LiF surface. As the

oxidation process continues, with mainly oxygen and water vapour diffusing through the LiF

to the metal surface, oxide eventually begins to form at the island covered areas. As those

areas begin to oxidize, the apparent thickness and rate o f oxidation increase. Henceforth, the

oxidation process follows inverse logarithmic growth as suggested by the Mott-Cabrera

exponential growth theory [6]. W ith a thinner LiF layer, this preferential oxidation ends

much earlier. For long exposure times, the estimated oxide thickness is the same as that o f

uncoated Mg, suggesting that the MgCCb dominated stage is much shorter than that with a

thicker LiF layer.

The surface structure would, therefore, not resemble the sandwich structure o f A1

oxidation. With this structure, and as a result o f porosity as the oxide layer forms [20], the

overlayer model would underestimate the thickness o f the oxide formed on the surface,

estimating a negative oxidation rate, as was observed in figure 7-4.

7.4 Chemical structure at organic interface with Mg/LiF in device structures

The oxidation behaviour observed at the metal surfaces exposed to air is magnified in a

device, especially since highly reactive Mg cathodes are known to cause molecular

fragmentation reactions at the interface with Alq3 (see chapter 5). With the presence o f LiF,

the oxidation at the interface would be expected to change, as was observed at

Al/LiF/organic interfaces. Previously, it had been proposed that a minor mechanism in

injection enhancement was the prevention o f interfacial reactions [21], However, this

assumption was based on Al/Alq3 interactions, where the reaction is inferred through small

modifications of the N Is core level peak. As was observed in chapter 6, the presence o f LiF



at the interface does in fact suppress interfacial oxidation for Al. Since the A1 is deposited on

top of the LiF, the logical assumption would be that the deposited LiF modifies the surface

reactivity o f Alq3, perhaps with the formation of charge transfer compounds (see Chapter 9).

With the surface reactivity modified or blocked, the organic surface would no longer be

susceptible to degradation.

With such an effect, the deposition o f reactive Mg on the LiF surface on Alq3 should

also show signs of reaction suppression, with less molecular breakdown. On the other hand,

if the reaction prevention has more to do with the metal/LiF interaction, then the impact at

the interface would be very different than that observed at Al/LiF interfaces with organic

molecules. As can be seen in Figure 7-6, the Al 2,p core level for the cathode side o f the

interface for Mg/LiF cathodes have the asymmetric line shape consistent with interfacial

breakdown [22], indicating that the LiF layer did not prevent the interfacial reaction between

Mg and Alq3. The composition ratios, in table 7-2, also support molecular fragmentation,

with the Mg/LiF surface showing a significant N deficiency, similar to that observed at Mg

surfaces. Since the theoretical N/Al ratio is 3 N per Al atom for Alq3, N deficiency is an

indication that the molecular structure is no longer consistent with Alq3.

Figure 7-6 Al 2p core level recorded for the Mg/LiF surface. The experimental data can be well fitted by the sum (solid line) of two separate peaks (dashed lines), one metallic state at 72.9eV and another Al3+ state at 74.4eV.

78 76 74 72 70 68

Binding Energy (eV)

Al 2p

CO

c

COEoz



Table 7-2 Atomic ratios at the cathode side of the as-peeled interface

Ratio Mg/LiFSurface

MgSurface

Mg2+/Mg 3.4 1.9

0/Mg2+ 2.5 2.0

C/Mg2+ 7.1 9.0

C 29O eV / M g 2 n d ox ide 1.0 0N/Al 1.3 1.1

Molecular breakdown reactions in Alq3 can be thought o f as metal exchange reactions

(see chapter 5), so the presence of metallic Al at these interfaces suggests that Mg should be

oxidized in both cases. Correspondingly, the Mg 2p core level for both Mg and Mg/LiF is

highly asymmetric, corresponding to multiple oxidized and metallic states. Figure 7-7(a)

indicates that the high binding energy peaks for the two surfaces are separated by about

0.5eV, indicative o f a different oxide phase forming at the interface with a LiF interlayer.

With alignment using the Alq3 3+ oxidation state for Al 2p at 74.4eV, the Mg 2p core

levels both show a metallic state at 48.5eV, slightly lower than that expected [12]. Figures 7-

7(b) and (c) show the curve fitting for the Mg 2p core levels at the interface with Mg and

with Mg/LiF. For the Mg surface, the Mg 2p core level is well described with only one high

binding energy peak, 1.4eV above the metallic peak, consistent with oxide and hydroxide

formation as on the exposed metal surfaces. The Mg 2p core level on the Mg/LiF surface

requires two peaks on the high binding energy side, at 2.1 and 3.5eV above the metallic,

suggesting the formation of a complex oxidation state, and perhaps some carbonate.

The F Is core level is unchanged and shows no evidence of a shoulder at 690eV, as

would be expected for F-C bonds [23], suggesting that the molecular breakdown reaction

supersedes the formation of charge transfer bonds described in chapter 8. The peak positions

and fitting conditions are listed in table 7-3.



Mg 2pC3

— • — M g / L i F — o — M g

JD

(0CD£(0c

CDcT3<DN15 «D£oz

Mg 2p

Mg 2p

i - v -

4854 52 50 46

Figure 7-7 Mg 2p core level for both Mg and Mg/LiF cathodes at the cathode side of the as-peeled interface, indicating (a) the difference in the high binding energy shoulder for the two surfaces of 0.5eV. (b) and (c) The curve fitting results of Mg 2p recorded on the Mg and Mg/LiF surface respectively. The experimental data (solid circles) can be well fitted by the sum (solid line) of separate peaks (dashed lines). In both cases, the metallic state is at 48.5eV. For (b) the oxide peak corresponds to hydroxide formation at 1.4eV above the metallic. The two peaks in (c) are located at 2.1 and 3.5eV above the metallic peak.

Binding Energy (eV)

CoreLevel Component BE (eV)

Mg

FWHM(eV) P* BE(eV)

Mg/LiF

FWHM P*

Mg 2p Metallic Mg 48.5 1.00 1.1 48.5 1.00 1.1

Mg “oxide” +1.4±0.1 1.82 - +2.1+0.1 1.72 -

MgC03 - - - +3.5+0.1 1.35 -

Al 2p Metallic Al 72.8§ 2.11§ 0 72.9 2.15 0

Alq3 74.4§ 1.75§ 0 74.4 1.66 0

* asymmetry parameter as defined in Appendix C (by p = TS*en ), §from chapter 5



The formation of different oxidation products at the organic/cathode interface with

different cathodes is consistent with the results o f oxidation of exposed Mg surfaces, where

the presence of LiF favoured the formation o f MgCC>3 over that of hydroxides. In the ion

exchange reaction between Al and Mg with molecular fragmentation, there would be an

abundance of carbon based fragments at the interface. The presence o f LiF may, therefore,

again be encouraging the formation of carbonate type structures. However, the higher

binding energy shoulder visible at the Mg/LiF surface is only 0.5eV higher than that at the

Mg surfaces, which is not consistent with MgCCb formation. The stoichiometic ratios ob

served at various surfaces do suggest that the molecular fragments may also be different for

Mg and Mg/LiF interfaces. Some stable potential molecular fragments incorporating Mg are

shown in figure 7-8, based on the observed atomic ratios from table 7-2 for the two cathodes.

Mg interface Mg/LiF interface

CMMg,

+

Figure 7-8 Potential reaction products formed at the cathode/Alq3 interface for Mg cathodes

The passivation of Al surfaces with LiF suggests that the coherence of an overlayer

may be having an effect on the activity o f metal surfaces. This is similar in concept to the

empirical determination of the protectiveness of an overlayer using the relative volume of the

oxide and the metal, referred to as the Pilling-Bedworth ratio [24], If an overlayer is not well

matched with the underlayer, it may not be able to form a protective barrier at the interface.

The distortion in the interface coherence can allow fast diffusion pathways for oxygen and

water, increasing the oxidation. The surface lattice constant, assuming ( lx l) surface



structure, therefore, could be used to predict which of these interfaces would provide a better

contact. In this case, the lattice matching can be defined from a coincidence-site lattice

concept [25], where the long and short axes o f the surface unit cell for a given plane are

matched to gauge the coherence of the interface. The misfit is defined by,

, a, - aRA = —-------------------------------------------------------(7-1)

aAwhere an is the surface lattice constant along a given direction on the surface plane.

The lattice misfit is given for Mg with a number o f overlayers in table 7.4 for a few

low index planes, showing the particularly poor matching o f Mg and MgCCb type structures,

as well as for LiF along some symmetrically equivalent planes. There are currently no

crystallographic data for the possible molecular reaction products at the metal/organic

interface. However, the average bulk lattice constant can be estimated using Girolami’s

method [26] to predict the density and by assuming a cubic, close packed structure for the

molecular crystal. Such an estimation indicates that there is very little difference between the

two lattice constants, but the estimated values (8.86A and 8.83A for Mg and Mg/LiF

respectively) are so large that ( lx l) coincidence is not possible. From table 7-4, it appears

that the reaction product at the Mg/LiF interface would be less well matched with Mg. Using

this approximation, however, both of the predicted molecules show relatively good matching

- (2x2) for {1000}//{111}, (2x4) for {1010}//(211), and (3x1) for (n02)//{001}. Better

estimation of the crystal packing, and confirmation of the reaction products using infrared or

Raman spectroscopy, would be beneficial in determining which reaction products were most

likely. The stoichiometry and relative position o f the Mg 2p peaks for the Mg/Alq3 interface

also suggests the formation of Mg(OH)2 as another potential reaction product for Mg/Alq3

interfaces. O f all the possible overlayers, it appears that Mg(OH)2 has the best matching for a

few low index planes, and is most likely to be the reaction product in the absence of LiF.



Table 7-4 Comparison of surface lattice constants with Mg along low index planes. LiF and the products of Mg oxidation have (lx l) coincidence along both a and c axes. For the molecular fragments, the smallest lattice misfit is given by ( lx l) for {1000}, (2x4) for {1010}, and (3x1) for (| io 2) planes of Mg.

Lattice Lattice LatticeMisfit A Misfit A Misfit A

({1000}/{1000}) ({1010}/{1010}) ( (l 102)/(l 102))

Mg/LiF§ 11.3% a/a^ll.3% ,c/a2=72.5

ai/ai=25.3%, a2/a2=l 1.9%

Mg/Mg(OH)2 1.9% a/a=1.9%,c/c=9.0%

ai/ai=1.9%,a2/a2=5.0%

Mg/MgC03 30.8% a/a=30.8%,c/c=65.3%

ai/ai=30.8%,a2/a2=55.2%

Mg/Mg-quin**5 2.4% 2a/a!=2.4%,4c/a2=4.9%

3ai/ai=8.0%,a2/a2=2.4%

Mg/Mg-benz**5 5.0% 2a/a=5.0%,4c/a2=7.4%

3a,/a1=10.4%,a2/a2=0.3%

Lattice constants from [27] and [16],*Mg/Alq3 interface product, * Mg/LiF/Alq3 interface product* Estimated based on close packed cubic structure assumption of molecular crystal from estimated density §{ 1000}, joiiojand (no2) symmetrically equivalent to {111}, (211), and {011} respectively [28],

Since deposited Alq3 films are smooth (see chapter 9), lattice matching between the

interlayer and the metal should contribute to the contact integrity. This suggests that contact

formation between Mg and Alq3 may be disrupted with the presence of LiF and the bulky by

products of interfacial molecular breakdown. This disruption of the interface could be one

explanation for the poor device performance with Mg/LiF cathodes compared to Mg, shown

in figure 7-9. Devices have high injection voltages, and complete suppression of luminance

within typical operating voltages with the introduction of a LiF interlayer. StoBel et al. [2]

also observed this decrease in device performance with introduction of a LiF layer, as shown

in figure 7-9(b). They claimed that the work function of the surface was increased with the

introduction of LiF, leading to the poor injection properties, unlike the behaviour observed

with any other cathode. From the simple analysis above, it appears that this performance

might be related to the change in the surface activity of the cathode, as evidenced by the

change in the surface oxidation.



— OALi F/ Mg —n— 10A LiF/Mg

0.8

s, °'6

0.2

□ □□ o - on n d n j

■ i I ■ t . i ■4 6 8 10 12

0.0

■2 0 2

T" ■ I'" ,■ I" ■ TO A LiF /M g

’ 10A LiF/M g^ 0.4

■ ...........................0 2 4 6 8 10 12 14 16

0.0

Voltage (V) Voltage (V)

Figure 7-9 (a) Luminance-voltage characteristics for Mg cathode devices with and without a 10A LiF interlayer (b) Cur rent-voltage characteristics adapted from M. Stofiel et al[ 2].

7.5 Summary

LiF overlayers have a significantly different impact on the oxidation of Mg surfaces

than was observed for Al, as described in chapter 6. Initially, there is preferential oxidation to

form MgCCh on the surface, with little apparent change in the oxide thickness. Due to the

poor lattice matching of Mg and LiF along symmetrically equal low index planes, the LiF

likely does not form matched overlayer islands on the metal surface. As oxidation continues,

oxygen and water diffuse through the LiF lattice and along the interface, and hydroxides

become the dominant components of the coating. When this occurs, the oxidation rate

increases rapidly, and the “oxide” thicknesses for the coated and uncoated surfaces become

similar. Irrespective of the thickness, the LiF coated surfaces show much greater proportion

o f MgCC>3, which show very poor lattice matching with Mg for low index planes.

These changes in the oxidation of the coated surface can explain the behaviour of

organic light emitting devices with reference to the interfacial oxidation. In a device with

much less oxygen available for reaction than for the multilayer structures exposed to air, the



effect of the modification of the oxidation kinetics with the introduction of an interlayer

should be greater. When the two materials are not well matched, such as LiF with Mg,

oxidation is not prevented at the metal surface. In addition, the different oxidation by

products catalyzed by the presence of LiF could in fact be exacerbating the degradation of

Mg based devices that incorporate a LiF layer. It would appear that the LiF layer does not

prevent interfacial reactions in systems such as Al/Alq3 by passivating the organic layer.

Rather the reactivity at the interface is driven by the interaction of the LiF layer with the

cathode, promoting the breakdown reaction with Mg/LiF cathodes, but protecting the metal

from oxidation in the case of Al/LiF.

Interlayers that are lattice matched over a broad range o f orientations should provide

better oxidation resistance in devices than poorly matching ones, as long as the interlayers do

not promote the formation of non-matching oxidation by-products at the interface. The bulk

lattice constants, therefore, may be used to predict the effectiveness o f the contact at cathode

interfaces in devices or o f a particular metal/interlayer combination as a cathode material.

7.6 References

1 M. G. Mason, C. W. Tang, L-S. Hung, P. Raychaudhuri, J. Madathil, D. J. Giesen, L.Yan, Q. T. Le, Y. Gao, S-T. Lee, L. S. Liao, L. F. Cheng, W. R. Salaneck, D. A. dos Santos, and J. L. Bredas, J. Appl. Phys. 89, 2756 (2001).

2 M. Stofiel, J. Staudigel, F. Steuber, J. Blassing, J. Simmerer, A. Winnacker, H. Neuner, D. Metzdorf, H.-H. Johannes, and W. Kowalsky, Synth. Met. 111-112, 19 (2000).

A. Turak, D. Grozea, Z.H. Lu, in preparation.

4 C. Chen, S. J. Splinter, T. Do, andN. S. McIntyre, Surf. Sci. 382, L652 (1997).


6 V. Fournier, P. Marcus, and I. Olefjord, Surf. Interface Anal. 34, 494 (2002).

7 T. Do, S. J. Splinter, C. Chen, andN. S. McIntyre, Surf. Sci. 387, 192 (1997).

8 J. van den Brand, W. G. Sloof, H. Terryn, and J. H.W. de Wit, Surf. Interface Anal. 36, 81 (2004).




10 J. E. Gray and B. Luan, J. Alloys Compounds 336, 88 (2002).


12 X. D. Peng and M D. Barteau, Surf. Sci. 224, 327 (1989).

13 S. Ardizzone, C. L. Bianchi, M. Fadoni, and B. Vercelli, Appl. Surf. Sci. 119, 253 (1997).

14 Fundamental XPS Data from Pure Elements, Pure Oxides XPS International Inc. (1999).

15 FactSage software and database, C.W. Bale, P. Chartrand, G. Eriksson, K. Hack, J. Melancon, A.D. Pelton, S. Petersen, W.T. Thompson, (2001).

16 P. Villiars and L. D. Calvert, Data for Intermetallic Phases (American Society for Metals, 1985), Vol. 3.

17 A. Lubezky, Y. Kozirovski, and M. Folman, J. Electron Spectrosc. Relat. Phenom. 95,37 (1998).

18 H. J. Freund, Faraday Disscus. 114, 1 (1999).

19 Y. Yuan, D. Grozea, S. Han, and Z. H. Lu, Appl. Phys. Lett. 85, 4959 (2004).

20 K. Asami and S. Ono, J. Electrochem. Soc. 147, 1408 (2000).

21 M. G. Mason, C. W. Tang, L-S. Hung, P. Raychaudhuri, J. Madathil, D. J. Giesen, L. Yan, Q. T. Le, Y. Gao, S-T. Lee, L. S. Liao, L. F. Cheng, W. R. Salaneck, D. A. dos Santos, and J. L. Bredas, J. Appl. Phys. 89, 2756 (2001).

22 A. Turak, D. Grozea, X.D. Feng, Z.H. Lu, H. Aziz, A.-M. Hor, Appl. Phys. Lett. 81,766 (2002).


24 N. B. Pilling and R. E. Bedworth, J. Institute o f Metals 29, 529 (1923).

25 A. Zur and T. C. McGill, J. Appl. Phys. 55, 378 (1984).

26 G. Girolami, J. Chem. Ed. 11, 962 (1994).27 W. B. Pearson, A Handbook o f Lattice Spacings and Structures o f Metals and Alloys

(Pergamon Press,, New York, 1967), Vol. 2.28 G. Bums and A. M. Glazer, Space Groups fo r Solid State Scientists, 2nd edition

(Academic Press Inc., Boston, MS, 1990).


Chapter 8

LiF interaction with organics

8.1 Chemical structure of Al/LiF/Alq3 in organic light-emitting diodes1

8.1.1 Introduction

Organic light-emitting diodes (OLEDs) have been intensively investigated in the past 20

years [1] from both an academic and industrial points of view, mainly due to their potential

applications in flat panel displays. The advantages o f organic devices include higher bright

ness, greater viewing angle, a wider selection of colors, lower voltage operation, easier and

lower cost deposition, and compatibility with flexible substrates.

1 First appeared in a slightly different format as Applied Physics Letters 81(17) 3173-3175, Copyright 2002, American Institute of Physics (reproduced with permission).

- 143 -


Chapter 8 Organic/LiF interaction 144

OLEDs have a multilayer structure composed o f a transparent anode, an organic

active layer, and a metallic cathode. Achieving improved device performance, such as higher

brightness and efficiency, requires optimization o f charge injection and transport. The

injection behavior o f the contacts strongly depends on the nature o f the electrode/organic

interface.

For the cathode/organic interface, enhanced electron injection is desired in order to

balance charge carriers in the active layer. Initially, this goal was pursued by utilizing low

work function metals and metal alloys, such as Mg, Ca, Li, or Mgo.gAgo.i [2,3,4,5], as

cathodes. Unfortunately, all o f these materials are highly sensitive to moisture and oxygen,

have high chemical reactivity, and are comprised o f fast diffusing species. Being much more

stable and resistant to oxidation, Al is a highly desired cathode material. However, it makes

a poor OLED cathode due to its high work function. In order to deal with this problem, a

thin interlayer was used at the metal/organic interface, which dramatically improved Al

performance as a cathode. This interlayer could be formed by depositing a thin layer o f LiF

[6], or other alkaline metal insulators [7]; by doping the near interface region of Alq3 through

coevaporation with Li [8]; or by doping the near interface region o f the Al cathode with LiF

or CsF [9],

Currently, the most widely used cathode is the bilayer Al/LiF (with ~ 5A thick LiF).

While the effect of LiF in improving device efficiency is well documented, the underlying

working principle is still not fully understood. A number of mechanisms explaining the

observed enhancement o f electron injection have been proposed, such as electron tunneling

through a thin insulator layer [6], band bending at the metal/organic interface [6], lowering of

the work function of Al [10], the presence of interfacial dipoles [11], and LiF dissociation

with released Li atoms reacting with Alq3 to form Alq" anions [7], Recent studies, especially



those showing similar OLED performance for LiF doped Alq3 cathode [12,13], perhaps rule

out the first three mechanisms. Whereas there seems to be a wide spread belief o f LiF

dissociation as the dominant mechanism [7,13,14,15], there is no conclusive direct evidence

for Li and F being in different chemical states than LiF [6,7,14,16] nor o f the strong ionic

bonding o f LiF being broken.

An understanding o f the structure and electronic properties o f the metal/organic

interface is further complicated by the fact that the interface is sometimes not abrupt,

extending for several nanometers [17]; that the interface is located deep inside the device;

and that the 5A thick LiF layer cannot form a continuous interlayer.

In the first half of this chapter, we report photoemission spectroscopic results on the

bilayer Al/LiF - Alq3 interface, based on a novel method o f exposing the buried interface

under vacuum conditions.

8.1.2 Experimental

The OLED structures were fabricated using an OLED cluster tool2 as described in chapter 4,

and have a Al/LiF/Alq3/ITO configuration with the thickness o f the LiF layer varying from

3 A to 20A. The samples for analysis were prepared by the peel-off method in vacuum. This

in-situ peel-off technique produces perfect cleavage at the metal/organic interface, which

converts the buried interface into an organic film surface and the Al/LiF surface. The XPS

spectra were generated by an Mg Ka source with photon energy of 1253.6 eV at a pass

energy o f 29.35 eV. For depth profiling analysis, the sputtering is performed using an 3 keV

Ar+ ion beam at 60° incidence angle.

2 Samples were fabricated at Luxell Technologies Inc. rather than using the cluster tool described in chapter 4; however, the methodology was the same.



8.1.3 Results and discussion

Figure 8-1 shows core level spectra o f Al 2p, O Is, C Is, and N Is measured at the organic

side of the interface. They are typical for all the samples examined, regardless o f the LiF

layer thickness o f 3A, 15A, or 200A. All the spectra were aligned consistently based on the

Al 2p peak position at 74.4 eV for the Alq3 compound [17]. The atomic concentration ratios

correspond to that calculated based on Alq3 molecular structure regardless of the interlayer

thickness as shown in table 8-1.

CO

' c13

_QL.

£w c0

~o0N

03E

(a)- O 1sA l 2 p

72 540 536 532 528

(c>C 1s ; n i s

288 284292 408 404 400 396

Binding Energy (eV)Figure 8-1 Various XPS core level spectra recorded on the organic side of the cleaved cathode/organic interface.

Table 8-1 XPS measured ratios on organic side of buried surfaces

LiF thickness O/Al N/Al C/Al

Alq3 3 3 27

3A 3.2 3.0 29.7

15A 3.0 3.0 27.2

200A 3.1 2.9 27.1



Previous XPS studies reported that LiF remained a stoichiometric compound and the

core level peaks showed no dissociation [6,7,14], However, Li has a photoionization cross-

section 40 times lower than that of F making core level analysis difficult. Moreover, all these

studies were inherently limited due to their simulation of an OLED structure through

deposition of a few monolayers o f metal on LiF/Alq3, instead o f investigating the

metal/organic interface in working devices. Another common feature o f these reports is the

shift to higher binding energy of the core level peaks of O Is, N Is, Al 2p, and C Is. In

addition, they observed a broadening o f O Is peak and a shoulder developing in the N Is

peak. The latter is believed to correlate with a reaction at the pyridyl ring o f Alq3. These

spectroscopic features were common to those reported during the deposition o f metals such

as Li, Mg, Ca, K, and Na directly on Alq3. Therefore, they were considered as indirect proof

of LiF dissociation and the presence of free Li atoms in the case of Al/LiF/Alq3.

In contradiction to these reports, we found that there is no peak splitting in the N Is

core level, as shown in Figure 8-1 (d). The current data indicate that the pyridyl ring is still

intact, consistent with atomic concentration data obtained on these samples. Nevertheless, F

species are detected on the organic side o f the interface, which is shown in Figure 8-2. Two

distinct F Is core level spectra were detected. The amount of F detected is small (between 0.3

to 0.4 at%), close to the detection limit of the instruments; however, ours is the first report of

F being found in a different bonding than LiF. For all LiF layer thickness cases, the positions

of the peaks were consistent and the split could be observed even for 3 A LiF(Figure 8-2 (a)).



Figure 8-2 F Is core level spectra recorded on the organic side of the cleaved interface with LiF interlayer thickness of: (a) 3A, (b) 15A, and (c) 200A, respectively. The curve fitting results are also shown for the 200A LiF case. The experimental data is well fitted by the sum (solid line) of two peaks (dashed line), one at 685.7 eV corresponding to a LiF bonding, and the other at 688.5 eV due to C - F bond.

Figure 8-2 (c) shows the curve fitting o f the experimental data obtained for 200A LiF

interlayer. Two peaks were identified; one at 685.7 eV corresponds to LiF bonding, and the

other at 688.5 eV is attributed to F attached through 7t-bond to C [18]. The latter peak

indicates the attachment o f the F to the C atoms from the Alq3 molecule and the presence of

an F doped layer in the interfacial region. This F Is shoulder was observed on the cathode

side of the interface as well, as shown in figure 8-3 for a 15A LiF layer, within the Alq3 layer

left on the surface. The shoulder disappears almost immediately upon sputtering, along with

the N Is from Alq3, indicating that the interaction is limited to a small interfacial region.

C

694 692 69 0 688 686 68 4

Binding Energy (eV)



F 1s

Figure 8-3 F Is core level spectra recorded on both the (a) organic and (b) cathode sides of the cleaved interface for LiF interlayer thickness of 15A.

03 FI s

c

692 690 688 686 684

Binding Energy (eV)

Possible sources o f F are the LiF deposition process or the dissociation of the LiF

layer at the interface with Alq3, but the intensity of the XPS signal from the expected free Li

atoms would be under the detection limit, preventing any conclusion on the Li chemical state.

Reported ultraviolet photoelectron spectroscopy (UPS) spectra [7,14,16] o f Al/LiF

thin layers deposited on Alq3 show a shift of the occupied molecular orbitals of Alq3 to

higher binding energy, which may lead to a reduction of the barrier height for electron

injection at the interface. This molecular orbital shift could now be explained by the

attachment of F to the n electrons on the conjugated ligand. F is likely attached to the non-

pyridyl side of the quinolate ligand, and gains charge from the molecule, due to its large

electronegativity. The loss of charge in the Alq3 molecule will lead to a reduction o f its

Coulombic potential, thus resulting in an increase in binding energy of valence shell

electrons, i.e., the molecular orbital features are shifted to higher binding energy.



Figure 8-4 displays the XPS depth profile results, from the 200A LiF layer specimen

for Al 2p, F Is, Li Is, O Is, and N Is core levels, on the cathode side of the interface.

Similar to the organic side, there is no additional shoulder or peak in the N Is spectrum. The

N to Al ratio is close to the value for Alq3, in contrast to N deficient ratio we reported

previously [19] for a Mg-Ag cathode surface. For that case, the low ratio o f N/Al indicated a

chemical reduction reaction between Mg and Alq3. Such a reaction did not take place for the

Al/LiF bilayer. The Al, O, and N on the cathode side of the cleaved interface is from Alq3

residual left on the cathode. Another important feature o f the depth profile spectra is the

diffusion of O from the exterior surface o f the Al cathode through pinholes and defects,

which results in the formation o f Al oxides. The presence of this additional Al oxide

enhances the complexity of the Al 2 p peaks observed for the 3 A and 15A LiF samples. The

depth profile also shows that the diffusion o f O ends abruptly at the Al/LiF interface. This

indicates that LiF interlayer acts as an excellent buffer layer limiting O diffusion from the top

surface into the organic film, as described in Chapter 6.

Chemical reactions between Al cathode and LiF layer, possibly facilitated by water

adsorbates, have been considered and proposed [15] for the formation of free Li metal atoms.

Our spectroscopic results from both side o f the interface do not show any reaction taking

place between F and Al, such as the formation of A1F3 (with the corresponding shifts in

binding energies o f Al and F) or other Al -F compounds. Additionally, alkali metal catalysts

have been used [20] to promote oxidation of Al surfaces. LiF does not dissociate on the

surface of Al, and so there is no enhancement of the surface activity. It has been established

that, in fact, LiF has the opposite effect at Al surfaces, passivating the surface against

oxidation. Since Al forms a metal excess oxide, if LiF were dissociating in contact with Alq3,

then the presence of Li+ ions at the interface should increase both the diffusivity o f metal



ions and the oxidation rate [21]. Therefore, one would expect to see enhanced oxidation and

molecular breakdown. Instead the opposite was observed, suggesting that LiF stays intact

during its interaction with Alq3.

F 1 s Li 1 sAl 2p

23 nm

x25

'(/)c0 -I—»

O 1 s N 1 s

c

Figure 8-4 XPS depth profile on the Al/LiF side of the cleaved interface with 200A LiF layer. The evolution of the Al 2p, F Is, Li Is, O Is, and N Is core level features is shown as a function of the

distance from the interface. The presence of an Al oxide at the Al/LiF interface suggests diffusion of O through pinholes in the Al films. Such O diffusion ends abruptly at the Al/LiF interface.

536 532 528 404 400 396

Binding Energy (eV)

8.1.4 Summary

Despite the fact that no XPS results show evidence of dissociation o f the ultrathin LiF layer,

earlier studies assumed this dissociation and the formation of free Li atoms, leading to radical

Alq3 anions. The present vacuum peel-off technique allowed the buried organic/cathode

interface to be investigated directly, which revealed the presence o f a new F-doped Alq3

layer instead of an abrupt junction between Alq3 and the Al/LiF bilayer. The attachment of F

to the conjugated Alq3 ligand may cause the shift of the molecular orbital levels to a higher

binding energy, leading to a reduction of interface charge injection barrier.



8.2 LiF interaction with C60

8 .2.1 Introduction

As described in the previous section, the interaction o f LiF with Alq3 is not similar to the

molecular breakdown reactions described in Chapter 5. The interaction in organic/LiF

systems appears to be dominated by a charge transfer between the conjugated C structure and

the LiF ions. As such, it cannot be described by the simple inorganic ionic state

approximation for organometallics. Metal-carbon based interactions in organic/inorganic

systems for OLEDs have been neglected for the large part in the experimental literature, as

MNDO and DFT theoretical calculations have shown that C-metal bonds tend to be less

energetically favourable than O-metal bonds [22,23]. Work done in our research group has

established that this organic-LiF interaction is related to the electronic nature o f the organic

layer, rather than a result o f functional group chemistry [24,25]. The formation o f the charge

transfer interaction during co-evaporation, visible by optical absorption and by XPS, is

limited to electron transport layers.

Correspondingly, improvements of the performance in OLEDs with a LiF interlayer

have been reported for other electron transporting molecules [26]. Flowever, the concept of

the LiF/metal combination as a universal cathode material is unfounded [27,28,29], In fact, it

appears unlikely that any particular cathode can be universally applicable for OLEDs, since

interfacial reactions can diminish potential improvements resulting from charge-transfer [30],

A key question lies in whether interfacial reactions follow a general predictable theory; or if

there are case specific descriptions for each molecule, such as the metal exchange reaction

description for Alq3.



Cm, an interesting optoelectronic material in its own right, is the ideal molecule to

study the interaction o f C with cathode materials. It represents an electron transporting

conductive molecular film composed entirely o f C atoms. The wealth of literature about C6o

has allowed the development o f a preliminary model o f this interaction to describe the

spectroscopic results. Although the spectroscopic effects are explainable for LiF-C6o, the

experimentation to date with other organic molecules has been somewhat inconclusive.

In the second half o f this chapter, we attempt to describe the experimental XPS

spectra using various theoretical models to account for the observed changes. Analysis o f the

best available information with these models suggests that the LiF-C6o interaction at the

interface is best described by the formation o f a charge-transfer bond. There has been some

indication that this C-F interaction may be related to the deposition conditions, and that,

contrary to the speculation in the first half o f this chapter, it is unlikely to be playing a major

role in device performance. Further work needs to be done utilizing other techniques to

clarify the growth mechanisms for LiF on metal surfaces and the impact that various

substrates have on the appearance of the charge-transfer complex.

8.2.2 Experimental

All the depositions were performed in the OMAC, and analyzed in-situ in O-MAC, for both

organic and LiF evaporation.

For deposition of organics on LiF, three types of substrates were used. The first,

referred to as LiF crystal, consisted of 1-2 mm thick rectangular sections cleaved from LiF

(100) single crystals, ranging between 5 to 10mm square. The second, referred to as

amorphous LiF, were 20mm square sections of 200 A thick LiF thermally evaporated onto



H-terminated Si (100) wafers. The third consisted of metal (Ag, Au, Pt) or ITO substrates

with a small amount (~5A) of LiF deposited prior to organic deposition. Organic molecules

were also deposited directly onto these underlying substrates. The metal films (Ag, Au, Pt) of

1000A were sputter deposited on Si using the Kurt J. Lesker cluster tool. The ITO film was

commercial grade ITO sputter deposited onto glass. For deposition o f LiF, the substrate was

350A of an organic molecule (Alq3 or C6o) thermally deposited onto Si.

Samples were produced using the deposition procedure as described in Chapter 4.

Briefly, organic molecules and LiF were thermally evaporated from crucible sources at an

average rate of lA/min and 2-3A/hr for organics and LiF respectively onto previously

prepared substrates. The growth o f C6o is described by the formation o f monolayers (ML)

rather than in A since the growth appears to be layer-by-layer. For Ceo, 1 ML is defined as

1.15xl014 molecules/cm2 for a close packed FCC structure, assuming a Cô diameter o f 7 A

[31] and a density o f 1.65g/cm3 [32],

8.2.3 Results and discussion

8 .2.3.1 F Is core level fo r C^j-LiF interaction

From section 8.1, it can be seen that the high binding energy shoulder in the F Is core level

may be attributed to a C-F interaction, related to the electron accepting ability o f the organic

layer [24,25], Therefore, as expected, the high binding energy shoulder in the F Is core level,

-3 .5 ± 0.4 eV above that of ionic LiF, was also observed after peel-off in device structures

where the Alq3 layer was replaced by C60, as shown in figure 8-5 below.



T J s

694 692 690 688 686 684

Figure 8-5 F Is core level spectrum of the cleaved interface of a single layer device with a glass/SiO/Al/LiF/C60/LiF/Al/SiO.

structure. Removal of the substrate and organic layers in vacuum left behind the cathode material (SiO/Al/LiF) and approximately 50A of C60 (Binding energy values not aligned externally).

Binding Energy (eV)

As C6o is both a good electron acceptor and has abundant C bonds, this shoulder also

appears during growth in-situ in the OMAC. The appearance o f this shoulder is independent

of the deposition order (figure 8-6).

Figure 8-6 F Is core level spectrum with high energy shoulder for (a) deposition of ~5A LiF on 350A C60 on Si (b) deposition of 10ML of Cm on 200A LiF on Si and (c) deposition of 2ML C60 on ~5A LiF on Au. The solid line in each case represents the LiF substrate, except for (a) where it represents crystalline LiF.

692 690 688 686 684Binding Energy (eV)



The observed shoulder is generally found with a binding energy around 688.5 eV.

Such binding energy values for F Is core level have also been observed in fluorinated

fullerene structures by others, and in those cases had been attributed to semi-ionic F

attachment to the molecule through the 7i-bonds [18]. As described in section 8.2.5 below,

however, the likelihood of LiF dissociation with F attachment to the Ceo structure is very

small. This suggests that the bonding between LiF and C6o is similar to the semi-ionic

bonding of F to C6o previously proposed for fluorinated fullerenes.

8 .2.3.2 Geometry optimized structures and theoretical prediction o f the F Is core level shift

To better understand the mechanism describe in the previous section, it would helpful

to determine the geometry optimized structures that result from the interaction o f C6o and

LiF, using density functional calculations. Such calculations were performed in collaboration

with Dr. Dharma-Wardana at the National Research Council. The Mullikan charges and

binding energies were determined for a number o f configurations for fullerenes interacting

with LiF molecules. The geometry optimized structures suggested that LiF molecules are

shortened, and transfer charge to the C6o molecule. The charge redistribution and shortening

of the bond would also be expected to affect the XPS spectrum by producing a high energy

shoulder, as was observed. However, the model focussed on individual molecule-molecule

interactions, neglecting relaxation effects. Therefore, the Hartree-Fock binding energy

values using Koopman’s theorem, given in Appendix E, underestimate the measured binding

energy o f solid state LiF by ~30eV. As there were also no features visible in the XPS

spectrum at those binding energies, it is likely that this covalent model description was too

simplistic to adequately describe the interaction between LiF and Cm- The electronic-



structure details, obtained from density functional calculations using the Gaussian-98 code

[33], are described in Appendix E.

Despite the problem of neglecting relaxation effects in the model, it cannot be

immediately dismissed. It is possible to describe the chemical shift, or the change in the

expected binding energy versus that o f crystalline LiF, since the Mullikan charges and bonds

lengths represent geometry optimized structures. To investigate such a possibility, the charge

potential theory of Siegbahn et al. [34], as described in chapter 3, may be used with the

estimated Mullikan charges. If the predicted shift is close to the observed value, this model

may still constitute a valid description o f the observed spectroscopic features.

Briefly, in the Seigbahn theory, a chemical shift due to a change in the local

environment can be expressed as from equation 3-5 by

AEb(F ls) = kjAqF + A V Mad + AE« (8-1)

where AEh is the change o f the core binding energy versus a reference compound, kj is the interaction coefficient between core electrons and valence electrons, Aqf is the difference in the effective local charge on the atom of interest, AVMad is the difference in the Madelung potential due to the surrounding atoms, and AEF is the difference in the relaxation energy due to photoelectron emission.

To correspond to the one molecule interaction described by the model, an ionic LiF

molecule can be taken as the reference compound, so that the effective local charge

difference can be estimated from the ideal ionic crystal and the predicted Mullikan charges.

The interaction coefficient, k, was taken as 19.6 for F Is, from the empirical formula

determined by Sleigh et al [35].

The Madelung potential using the shortened bond lengths can be approximately

described by the Coulomb interaction given in equation 3-6 assuming each atom as a point



charge in space (in this treatment, C6o is taken as a single point charge of molecular radius).

For crystalline LiF, the point-charge model can be described by

yMad = 25 .166V — = 25.166— (8-2)t j ru a 0

where qj is the charge on the other atoms, rtj the interatomic spacing to the atom of interest, a0

the crystalline lattice constant, all in atomic units. For LiF, the Madelung constant, ocuf is 1.75 [36],

For the individual molecules used in the model, however, all o f the nearest neighbour

interactions must be calculated, since the Madelung constant is simply taken as 1 in the

absence o f long range crystal structure.

Assuming that equal shifts in binding energy in all core levels occur with a change in

the chemical environment, the relaxation energy neglected in the model can be estimated

from the change in the modified Auger parameter, as defined by

Aa'= 2AE« = E b{F h) + E k(FKLL) (8-3)

where Ev(Fkt t ) is the kinetic energy of the first peak of the Auger transition associated with excitation in the F Is core level.

In general, the F Is shoulder only appears near the limits o f XPS resolution for F (0.3-

0.4at%). As such, it is difficult to observe the lower intensity F k l l Auger transition.

However, where both the F Kll and the shoulder on the F Is core level are visible, the F Kll

peak appears to have only one edge. As there are two peaks in the F Is core level, the

difference between them is also the difference in the Auger parameter; however, the value o f

the Auger parameter for the peak associated with ionic LiF is approximately leV higher than

the expected value for LiF thin films on surfaces [37], Therefore, the difference in the

relaxation from the reference compound, ionic LiF, is a combination of these two differences

in the Auger parameter. One can assume that the same relaxation is occurring in all o f the



covalent LiF structures. The predicted binding energy shift from only the initial effects can

be taken as approximately 0.5eV, rather than the 3.5eV that appears on the spectrum.

Table 8-2 lists the predicted simplified Madelung energy and binding energy shifts

from the Mullikan charges and bond lengths predicted by the theoretical model. The LiF-C6o

interaction appears to be consistent with the observed binding energy shifts, especially for

configurations with Li pointing toward the C6o molecule. From this analysis, it appears that

the configuration with one LiF molecule oriented normal to a hexagonal face, with Li

pointing towards that face, best predicts the shift in the F Is spectrum.

Table 8-2 Theoretical binding energy shifts for model structure of LiF-C60 interaction

Structure <lu qcJ c-60 r FLi (A) r FCm (A) VF (eV)

LlFCryst -1 +1 - 2.01* - 12.52 -

T iFmonomer -1 +1 - 1.51s - 9.53 -3.0

LiF -0.5 +0.5 - 1.586 - 4.54 1.6

Cso-FLi1 -0.491 +0.523 -0.0365 1.573 4.5* 4.71 2.0

C60-LiF2 -0.5 +0.365 +0.135 1.574 6.6* 3.63 0.71

C60-LiF3 -0.499 +0.348 +0.153 1.577 4.5* 3.67 0.24

LiF-C6CrLiF4 -0.493 +0.535 +0.042 1.573 4.5* 5.47 2.6

-0.504 +0.347 +0.157 1.574 6.6* 3.55 0.55

LiF-C60 -FLi5 -0.490 +0.524 +0.034 1.573 4.5* 4.94 2.2

-0.489 +0.521 +0.032 1.573 4.5* 4.91 2.2

FLi-C60-LiF6 -0.480 +0.374 +0.106 1.572 6.6* 3.85 1.3

-0.489 +0.355 +0.134 1.573 6.6* 3.72 1.2

From a„ for crystalline LiF after Euwevna et a/.[38]. From monomer size during vapour deposition [39] From C6o diameter after Troullier et al. [31], and assuming the LiF molecule is lA away from the surface o f Cso.1F near a hexagonal face, the LiF bond is normal to the face.2Li near a hexagonal face, the LiF bond is normal to the face.3Li is on a bond between a hexagonal and a pentagonal face. The LiF bond is slanted so that the F atom is

above the pentagonal face.4 F near a hexagonal face, Li near the opposite hexagonal face. LiF bonds normal to the hexagonal faces5 same as above but an F is adjeacent to a hexagonal and its opposite hexagonal face as well6 same as above but with Li adjacent to both hexagonal faces



However, it is difficult to definitively ascribe the shift to any configuration for the

LiF-C6o interaction, as the differences in the predicted chemical shift for different structures

are well within the error in the relaxation energy. Since the Auger peak is not as sharp as the

core level, the binding energy is only accurate to within ±0.5eV, and the error in the Auger

parameter is nearly leV. This is further complicated by the fact that the selection o f a

particular binding energy for charging alignment for an insulating compound like LiF effects

the value of the Auger parameter; increasing the expected error even further [40]. Only those

cases were F is pointed towards the C6o molecule would likely fall outside of this range, and

so are eliminated as possible configurations.

From this analysis o f the relative binding energies, the covalent model cannot

completely be eliminated as a possible description o f the observed spectroscopic results,

suggesting that LiF molecules are oriented with the Li pointing towards the molecule. During

thermal evaporation, LiF vapour is known to consist of monomers, dimers and trimers

[41,42]. The experimentally determined monomer length, 1.51 A [40] is very similar to the

predicted value for both isolated LiF molecules, and the LiF interacting with C6o- For LiF

deposition on the organic surface, therefore, some of the LiF could be interacting with Ceo,

without relaxing to the length expected for crystalline LiF.

However, the appearance of this shoulder during deposition of C6o on LiF makes the

covalent model of the LiF-C6o interaction fairly unlikely. As described below in section

8.2.3.4, the high cohesive energy of LiF suggests there is a high driving force for LiF cluster

formation. Additionally, deposited polycrystalline or amorphous films would still be

expected to have Voroni cells approximately the size o f the crystalline lattice [43]. For thick

films of LiF and for the LiF crystal, there is unlikely to be any such covalent LiF monomers.



The observation o f the high binding energy shoulder with C6o deposition on such surfaces,

and the unrealistic binding energy values from the theoretical model, suggest that covalent

LiF formation cannot explain the spectroscopic results, even though the observed chemical

shift may be similar. A more realistic model o f the interaction between LiF and C6o would

need to use a molecular cluster or a jellium type approach for LiF [44] to account for the

relaxation energy.

8 .2.3.3 C Is shake-up satellites fo r deposited monolayers

Since the covalent model is not likely, the best description is still that o f a charge transfer

bond between Cgo and LiF. Although the impact of the LiF-C6o interaction on the C Is

spectrum is quite subtle compared to that observed on F Is, there is evidence o f some charge

transfer interactions. There were, however, no observed shifts in the main C Is core level, as

would be expected if covalent C-F bonds were forming. Since C6o has an extensive shake-up

structure on the high binding energy side of the C Is core level, the observed peak at 290eV

cannot be definitively assigned to F-substituted carbon alone. If both C6o and C-F bonds were

contributing to the spectrum, however, the intensity of the 290eV feature could have been

higher than that of C6o alone. In our study, the opposite was observed, with the satellite

intensity smaller than that expected, suggesting that F is not covalently bonded to C.

The attenuation of the high binding energy satellites and change in the relative

contribution from the various shake-up transitions indicates a change in the local distribution

o f delocalised electrons due to the presence of LiF, regardless of the deposition direction, as

shown in figure 8-7.



^ _□ —350A on Si

. —o— 5A LiF on 350A i- Q —A— 6 ML Cffi on 5A LiF

( 0 - v - 2 M L C K on5ALiF _ ■ theoretical shakeu i structure

ifwM. |Figure 8-7 C Is high binding energy satellites. The dropdown lines indicate the theoretical position of the shake-up features after Enkquist et al. [45]. The oval indicates the missing p-p* feature for LiF-C60, observed for pure deposited C60.

9 8 7 6 5 4 3 2 1Binding Energy (eV)

From figure 8-7, it can be seen that the major change in the satellite features occurs at

around +6eV from the position o f the main C Is core level3. This theoretically predicted [45]

and experimentally [32] observed feature for C6o corresponds to the superposition of a k-k*

dipole shake-up and a broad n plasmon [45]. As the amount of deposited C6o was increased,

layer of LiF, this suppressed feature begins to emerge, as shown by the upward triangles in

Figure 8-7.

Shake-up satellites are reflective o f the valence band characteristics, since they are

based on elastic loss processes within the orbital structure of the atom, as described in

Chapter 3. These satellites can be described as excited states of the molecule due to the

change in the potential of the molecule with the creation of a Is core-hole [46,47]. The

energies of these excited states after the loss o f a photoelectron can be approximated by the

ground-state energies of the molecule with promotion of a valence electron to an available

3 Note that satellite peak positions are referenced to the position o f the mainline at 285eV. This convention is used throughout this discussion.

and the C-F shoulder in the F Is spectrum was no longer visible at 6 ML deposition for a thin



empty state. The relatively low resolution of XPS generally precludes the use o f the satellite

structure to draw conclusions about changes in the valence band. However, C6o has a

particularly well differentiated and well described [32,45,48] set o f satellites, which aids in

spectral interpretation. A high resolution scan of the shake-up features for a thick layer o f C6o

is shown in figure 8-8, where all o f the theoretically predicted peaks are clearly visible. The

orbital structure of C6o indicating the observed transitions are listed in figure 8-9 below.

Almost all o f the shake-up structures correspond to some transition to the 5tiu*

(LUMO) level, giving a clear picture of the valence band characteristics o f C6o- The shake-up

feature at 1.9eV from the mainline o f the core level corresponds to electron promotion from

HOMO to LUMO. The sharp features at 3.8 and 4.8eV correspond to a number o f monopole

and dipole transitions, and finally the last sharp feature, at 6.0eV, corresponds to a

superposition of a n-K* dipole transition from a low lying orbital (6hg) to the LUMO and a

broad n plasmon [45],

C 1 s

8 6 4 3 2 17 5Binding Energy (eV)

Figure 8-8 High resolution scan of satellite structure for C60. Drop down lines represent the theoretically determined orbital transitions from Enkvist et al. [45] all visible in the spectrum.



2 - .

0 H->

E>C -2

LU O)

. £ - 4■OcCQ

- 6 -

59u8h<2t„

25tig*5tI* LUMO

5hu HOMO

7h +5gn

4h + 4 t +6hu 2u g

h9

t„

- 8 -

Figure 8-9 Molecular energy levels of C6o (neglecting core hole ionization) (after [49 and 45]). The transitions that correspond to the observed features in the spectrum are the HOMO-LUMO transition between 5hu and 5tlu* at 1.9eV, and the dipole transition from 6hg to the LUMO at 6.0eV. The features at 3.8 and 4.8eVcannot be assigned to a single transition, but represent the (5hu, 7hg, and 4gg ) -> (5tiu*, 5t2u*, 8hg*, and 5gu*) monopole and dipole transitions.

For the LiF/C6o system, the attenuation o f the satellite intensity can be used as an

indication of chemical bonding from mixing of the C6o frontier orbitals. Purely physisorbed

molecules show satellite structures that replicate their gas phase analogs [47]. The

suppression of the satellite features is very common for C6o deposition. For C6o growth on

Au, which interact strongly [46], the effect on the satellites is very similar, requiring 5ML for

the satellite features to be truly resolved, as shown in figure 8-10. However, the nature o f the

substrate has an effect on the emergence of these satellite features, as the 6.0eV feature is not

evident if the high binding energy shoulder o f F Is is visible. For a 200A thick LiF

substrates, the satellite features for C6o without any interaction did not emerge even after

10ML deposition, where the F Is shoulder was still visible. By comparison, this feature was



already visible after 3ML when C6o was deposited onto a bare ITO substrate. ITO had been

chosen as an inert substrate since In does not tend to bond with C6o [46]. The fact that the

satellite features appeared only after 3ML deposition suggests that ITO is not an inert

substrate for C6o, likely due to the presence of O4.

‘ I I(a)C 1s C 1s

6 ML,

5 ML/

3 ML

1 ML

2 1 9 8 7 6 5 4 3 2 1

Binding energy (eV)Figure 8-10 C Is shake-up structure for C60 on (a) pure Au (b) 5A LiF coated Au

The LiF/C6o interaction is very similar to others that show chemisorption and slight

charge transfer. Enkvist et al. [45] attributed the shake-up feature at 6.0eV to the p bond on

the six membered rings of C6o- The suppression of this feature of the satellites with LiF

interaction suggests that the presence of LiF disrupts the shake-up of electrons involved in

van der Waals bonding between C6o molecules. Other systems with C6o absorbed onto a

substrate [46] have shown a similar attenuation and loss o f differentiation o f the satellite

structure with chemisorption, as with Au and Cr, and no loss of differentiation with

physisorption onto GaAs. This loss of the satellites during chemisorption is due to the

electronic coupling of an adsorbate hole and substrate excitations via relaxation processes

4 See section 8.3.2.5 for more discussion on this observed effect.



[47]. The attenuation of the 6 eV feature in particular can be explained using the theoretical

model of Schonhammer and Gunnarsson for satellites in chemisorbed adsorbate structures

[50]. In that description, a shake-up channel for the adsorbate is blocked due to a transfer o f

charge from the substrate to a previously unoccupied adsorbate valence level. In strong

bonding, the probability o f this charge transfer screening taking place is high and the satellite

feature will be greatly attenuated. Since the k -tz* transition represents a bonding-antibonding

band excitation into the LUMO, chemisorption between the adsorbate and the substrate is

shown by the suppression of this feature. The strength o f the bonding cannot be

quantitatively established from the amount o f attenuation because the presence of the bulk p

plasmon superimposed on the dipole transition prevents the feature from completely

disappearing from the spectrum. The features at 3.8eV and 4.8eV represent multiple

transitions, some to higher levels, which are less likely to be filled during charge transfer;

therefore, these excitations are still expected.

8 .2.3.4 Theoretical support fo r LiF "C m complex formation

The spectroscopic results appear to indicate that there is some sort of chemical bonding

interaction between LiF and C6o, consistent with a charge-transfer compound and no

dissociation of the LiF molecule. There is some controversy in the literature about the nature

o f LiF-organic molecule interactions, with some investigators claiming that performance

improvements occur due to LiF dissociation in contact with the organic molecule and a metal

[51]. All of the studies that claim dissociation did not mention the state of F in these systems,

offering only indirect proof. Subsequent investigations have shown no evidence of AIF3

formation, which would be expected with dissociation [52],



To eliminate the possibility o f carbon fluorination with LiF dissociation, the

following model reactions may be used to determine if dissociation is thermodynamically

viable.

LiF + Cgraphile => CF + Li (8-4)

LiF + C 6 H 6 => C6 H 5F + LiH (8-5)

xLiF + C60 => CmFx + xLi (8 -6 )

From these, the Gibb’s free energy change can be estimated to determine the

likelihood of carbon fluorination due to LiF dissociation.

Although no thermochemical data exist on low order fluorination of C6o, it can be

extrapolated from the existing data on high order C6oFx. The enthalpy o f formation has been

observed to have a linear relationship with C/F ratio for fluorinated graphitic structures,

independent of the type o f fluorinated C material [53], as also mentioned by Papina et al. [54]

for fluorinated Cgo Similarly, the entropy of formation for the stable forms of fluorinated C6o

show a downward trend [55], Therefore, using the data of Papina et al. for C60F36 and C60F48,

the Gibb’s free energy of formation for low order C6oFx can be estimated. Due to bond

equalization requirements to maintain a cage-like structure, ultralow fluorination is not

thermodynamically favourable. Therefore, C60F2 is the smallest stable fluorinated C60 [56].

Table 8-3 Gibb’s free energy of fluorination reaction at 298 K

LiF Â 1F3* LiF ->CF LiF C6H5F LiF->C6oF36 L iF ^ C 60F2

A G°rxn

(kJ/mol)541.9 808.3 391.4 1557.2 356.73

*assum e A1F3 formation in m etal exchange reaction with A lq3. T herm ochem ical data from T herm ochem ical H andbookfor LiF and organics, C 60 from K olosev [57 ], C 60F36 from Papina et. al. [54]



From table 8-3 above, it appears that a reaction between LiF and C would be

thermodynamically unfavourable. Even with low fluorination, LiF dissociation and Li+

doping of Ceo is unlikely.

Although carbide formation through dissociation appears to be unfavourable for the

LiF/C6o system, formation o f an intercalation compound may be possible, with diffusion o f

ions through the C6o structure. Wertheim and Buchanan [58] developed an empirical model

to predict the potential for bulk intercalation of C6o with various metals related to the

cohesive energy. They claim that if the empirical “intercalation” energy is positive, bulk

intercalation can be expected. Since crystalline LiF has a high cohesive energy, 10.7eV [59],

it is definitely not expected to have any intercalation effects [Einr=-17.5eV]. However, during

evaporation at room temperature, it is likely that crystalline clusters can form on the surface,

seen by the formation of LiF islands when small amounts o f LiF are deposited on metal

surfaces. These clusters would have a lower cohesive energy than in the bulk. Following the

models o f Qi et al. [60] and Pacchioni et al. [61] for FCC metals, the lowest limit o f a 6

particle cluster predicts an energy approximately half that o f the bulk cohesive energy value.

In these models, the metallic bonds were approximated as nearest-neighbour ionic bonds. As

LiF has an FCC Bravais lattice, this model can also be applied to the ionic bonds in LiF; even

though it is not a metal, LiF should fall under the theoretical description as it is made up of

ionic bonds. With this minimum value of half the cohesive energy, however, intercalation

still does not appear favourable for C6o/LiF. On the other hand, the model of Wertheim and

Buchanan does not apply for systems that form interfacial charge transfer compounds with an

intact inorganic layer. In fact, the inability to form intercalation compounds in this way does

not preclude the possibility of charge transfer reactions [46,58], only that the bulk formation

of such a compound is unlikely due to the precipitation of LiF clusters. Other systems, such



as Al, with an intercalation energy similar to that estimated for LiF, shows a tendency to

form dilute solid solutions with C6o even at room temperature [62].

The spectroscopic results, therefore, cannot be attributed to either the formation o f

carbide compounds or bulk intercalation through diffusion of ions. A likely explanation then

is the transfer of charge from LiF to C6o without dissociation in the formation o f a LiF#*C6o

complex. Cô shows a tendency to form such charge-transfer complexes near interfaces with

a variety o f metal systems [46,58]. It is unusual to observe inorganic compound charge

transfer complexes o f this nature since the electrons should be tightly bound; the adsorption

of CO on LiF, the closest model system to Ceo-LiF adsorption, for example, appears to be

totally electrostatic in nature [63]. However, since Li is a congener to H, the charge transfer

bond can be thought to take the form o f so-called lithium bonding, and the resulting complex

can then be considered analogous to H bonding complexes. Ammal and Venuvanalingiam’s

theoretical study [64] o f p-bonded systems as a bond acceptor from LiF indicates that the

lithium bonded structures show an interaction between the 7tc=c orbitals and the antibonding

a* orbital in LiF, with LiF complexes proportionally stronger than Li atom complexes. This

bonding would be disrupted by the presence o f Li+ ions from dissociation, as ionized

complexes would be stronger than those from LiF, again indicating the likelihood of intact

LiF molecules in contact with C(,o.

8 .2.3.5 Growth morphology and critical thickness fo r F Is peak appearance

There appears to be a critical thickness of C6o during deposition on LiF before the F Is

shoulder becomes visible, varying with the amount o f LiF available for the interaction. The

growth evolution of the high energy shoulder is shown in figure 8-11. As can be seen in



figure 8-11(c), even with only ~5A of LiF on the Au substrate, 2 ML o f C6o were necessary

to resolve the second peak. For 200A LiF, more than 5 ML were necessary before the peak

was resolvable. This is contrary to the expected interaction o f C60 and a substrate that shows

charge transfer, which generally only occurs in the first monolayer [46].

One possible explanation is the formation of clusters o f LiF on the Au surface during

deposition, whereby C6o in the first monolayer would interact with the substrate

preferentially to LiF. Subsequent deposition of C6o would then begin to show interaction with

the LiF clusters after the bare Au surfaces were completely covered. With the greater

thicknesses of LiF, which likely form a complete layer, the signal from the small amount of

C6o interacting with LiF is not visible spectroscopically due to washout from the strong signal

o f the substrate. This could support the fact that the LiF-C6o interaction was weak compared

to LiF cohesion, as expected. With sufficient C6o deposited to attenuate the substrate signal,

the bonding becomes visible. Two other substrates, LiF on Ag and ITO as in figure 8-11(a)

and (b), show no shoulder even after 2ML of deposition. These substrates were initially

chosen to mitigate the substrate-C6o interaction, but appeared to have had the opposite effect.

This could be a further indication o f large LiF cluster formation on the surface o f the

substrate, with the C6o interacting preferentially with the substrate in those cases.

The shake-up satellites for C Is for 2ML of C6o deposition onto LiF coated surfaces

in all cases show the suppression of the high binding energy satellite associated with the

formation of the charge transfer complex, though the structures are varied for different cases

(figure 8-12). Since the substrate properties have a major impact on the satellite structure

[47], this indicates that the interaction is strong in all cases, but different for different

substrates. LiF coated Ag in particular has a structure that can be resolved into two peaks,

rather than three as is usual for CLo-


Nor

mal

ized

In

tens

ity

(arb

. un

its)


w'c

-Qv.3*

'c/5c0c"D0N

I d£oz

F 1 s

2 ML

F 1 s

2 ML

F 1 s

T l s8 ML /V

LiF crystal

Figure 8-11 Evolution of the F Is core level with C&) deposition on a variety of substrates, (a) ~5A LiF on Ag (b) ~5A LiF on ITO (c) ~5A LiF on Au (d) 200A LiF on Si. Each cycle represents roughly 1 monolayer deposition of C60.

692 6 9 0 6 8 8 6 8 6 6 8 4 682

Binding Energy (eV)

ITO' — o — 5A LiF on Au

- a - 5A LIF on ITO —v — 5A LiF on

- c 60theoretical shakeui structure A

J>4 ^

o Va *

li7 6 5 4 3

Binding Energy (eV)

Figure 8-12 C Is shake-up satellites for 2ML deposition of C60 on a variety of substrates.



Some more work needs to be done using STM and AFM to examine the growth mode

of LiF on a variety o f surfaces and determine if they exhibit different sticking coefficients.

LiF deposited on the Ag substrate (figure 8-11) also appears to have a different chemical

state than that expected for LiF, with a low binding energy shoulder on the F Is spectrum. As

an in-depth study o f LiF growth is beyond the scope o f this project, some further work needs

to be done to confirm the activity of the surface by looking at growth of LiF on a few metal

surfaces of importance for OLEDs, such as Ag, Mg, Pt, Au, and Al. This would aid in the

establishment o f a physical picture of the interface formation process. Initial results indicate

that the deposition rate and chamber partial pressure may be playing a major role in the

appearance o f this shoulder in the F Is core level.

8 .2.3 . 6 Other LiF/organic interactions

We have also examined the interaction between LiF and C in a few other organic/LiF

systems in this study. The high binding energy shoulder in the F Is spectra was intermittently

observed, in addition to the C6o-LiF results outlined above, as summarized in Appendix F.

Though a wide variety of organic molecules have been examined in a number o f different

situations, the results remain inconclusive. It is interesting to note that though Alq3-LiF

combinations showed the F Is shoulder during co-evaporation and in devices, a similar series

o f depositions o f Alq3 on LiF substrates or LiF on Alq3 as for C6o did not give rise to the high

energy shoulder in the F Is peak. This could indicate that the activation energy for this bond

formation is much higher for Alq3 than for Cô, as would be expected since Alq3 has many

fewer 7t electrons per molecule at an equivalent thickness. As well, steric considerations

could be affecting the availability of the 7t electrons for bonding in a monolayer, since the

layer consists primarily of meridinal Alq3.



Although the results are inconclusive regarding the C-F interaction, there are a three

effects with other metal/LiF/organic systems that have been observed, which are reflective o f

our current understanding of this phenomena.

8.2.3.6.1 Effect o f interfacial reaction

The high energy shoulder was often observed during peel-off o f devices, for both Alq3 and

C6o electron transport layers. The only time this was not observed in a device was when there

was a molecular breakdown reaction between the cathode and the organic layer, such as for

Mg/LiF/Alq3 based devices as described in chapter 7. Another example is with a triphenyl

triazanine (TPT) electron transport layer, which appears to react with LiF and breakdown, as

shown by the N Is core level, as in figure 8-13 below. This suggests that a molecule likely to

react destructively does not form a charge transfer complex.

Cathode side

- * - 5 A LiF/AI ' A 3A LiF/AI

— A —TPT/5A LiF/AI- - - TPT/AI

Organic side

406 404 402 400 398 396

-5 A LiF/AI -TPT/5A LiF/AI - 3A LiF/AI

F 1s

cts

55c0c

T30N

696 692 688 684 680

Binding Energy Binding EnergyFigure 8-13 Normalized N Is core level for the cathode side of the cathode/organic interface for Al and LiF/AI cathodes, with TPT and Alq3 as the electron transport layer. The TPT/AI (closed triangles) and AlqVcathode show peaks consistent with TPT powder and the organic side of the interface, respectively.



8.2.3.6.2 Effect o f changing the cathode material

Initially, as described in section 8.1, it had been speculated that the C-F interaction

observed in Alq3-LiF devices might explain the performance o f devices with LiF interlayers,

since it appears in cases with improved performance such as Al, but is absent in with the use

o f Mg, which has much worse device performance. However, this interaction was also

observed for Ag/LiF cathodes, as shown in figure 8-14 below. LiF provides no performance

improvement for Ag based cathodes with Alq3 [65], It is likely, therefore, that though this F-

C bond is a spectroscopically observable phenomenon, and may be beneficial in improving

contact adhesion, it likely has little real impact on the device properties.

I 1------- 1------- 1------- 1

F 1sAl/LiF cathode side J

—• — Al/LiF organic sidesj

%

I 1 I 1 I Ag/LiF cathode side

— Ag/LiF organic side

R

694 692 690

Binding688 686 684 682

Energy (eV)

Figure 8-14 F Is core level for the cleaved surface, both the cathode and organic sides, for (a) Al/LiF cathodes and (b) Ag/LiF cathodes.

8.2.3.6.3 Effect o f limitations o f experimental set-up

In general, it has been difficult to observe the high binding energy shoulder during in-

si tu deposition, whether depositing organic molecules on LiF coated surfaces or LiF on

organic surfaces. The inevitable contamination of the OMAC chamber appears to be playing



a role in the intermittent appearance o f this peak. It was even observed during the initial LiF

deposition onto the metal surfaces. The appearance o f this shoulder was always accompanied

by the deposition of C on the surface, suggesting that the spectroscopic feature at 690eV can

likely still be attributed to a C-F interaction. One possible explanation is the interaction o f the

LiF molecules during evaporation with volatile organics within the vacuum chamber.

One of the limitations of the current experimental setup has been the low deposition

rates achievable with the cathode source. Even at the maximum output, the deposition rates

for LiF never exceeded 2-3A/hr. The likelihood o f incorporation of impurities within the

growing film for such slow growth is quite high, even in relatively high vacuum. As organic

evaporation was also carried out in this chamber, the potential desorption o f molecular

fragments from the chamber walls was possible. Chamber walls would sometimes reach

temperatures as high as 60-80°C during LiF deposition. Although the pressure did not

increase substantially during deposition, it is speculated that the chamber walls may have

been outgassing molecular fragments.

These fragments could have become incorporated into the molecular beam of LiF,

ultimately depositing as a combination o f pure LiF and C-LiF. When such a layer was

formed on the surface, the likelihood o f LiF interacting with the organic molecules

subsequently deposited on the surface diminishes. A good example o f this effect was

observed in C6o films grown on LiF coated Pt. C6o is known to form strong covalent bonds on

the surface of Pt [6 6 ], This was confirmed by us with the suppression of the satellite features,

also observed in the EELS spectrum by Cepek et al [6 6 ], When C6o was deposited on LiF

coated surfaces, the satellite features were resolvable much sooner. This suggests that the

strong chemisorption of C6o on Pt was disrupted, not by LiF, but by the carbon contamination



layer bonded to the LiF on the surface. Further investigations need to be done to find an inert

surface for C6o to confirm this effect.

Though current manufacturing practices call for much higher deposition rates than

were achievable here, the slightly lower vacuum that generally exists during device

fabrication may explain why this F-C interaction was observed during peel-off o f devices.

Recent evidence has suggested that ultra-high vacuum environments can suppress device

performance [67,68], since interfacial oxides are not readily formed.

8.2.4 Summary

Analysis o f the best available information with a few theoretical and empirical

models for C(,o and LiF suggests that the LiF-C6o interaction at the interface can be best

described by the formation of a charge transfer bond, regardless o f the deposition direction.

The appearance o f the high binding energy shoulder in the F Is was accompanied by a

change in the C Is satellites for C60 similar to what has been observed previously for

chemisorption type interactions. However, the interaction between LiF and organic

molecules generally is complex, and the simple models proposed here are not sufficient to

draw any new conclusions about the interaction. Additionally, the dependence of the

appearance of both the shoulder in the F Is core level and the satellite features on the type of

substrate used, the thickness of the deposited amount o f material, and the deposition

conditions requires further study. Further work needs to be done utilizing other techniques to

clarify the growth o f LiF on metal surfaces and the impact that various substrates have on the

appearance of the charge-transfer complex.

Initially, it had been speculated that the C-F interaction observed in Alq3-LiF devices

might be a reason for the improved performance with the use o f LiF interlayers. However,



the appearance of this interaction for Ag/LiF cathodes calls this speculation into question. It

is likely, therefore, that though the F-C bond observed for Alq3 and C6o molecules is a

spectroscopically observable phenomenon, and may be beneficial in improving contact

adhesion, it likely has little real impact on the device properties. Further work may be done to

examine the impact o f an ultra-clean environment and o f faster deposition rates on the

appearance of this high energy shoulder.

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Chapter 9

LiF layer properties

9.1 Introduction

Multilayered cathodes are widely used in organic electronics to promote electron injection, as

the device properties are largely controlled by the nature o f the organic electrode interface.

Generally, the interlayers are materials with high dielectric constants, such as LiF, which are

essentially electronic insulators. As such, their use is usually limited to ultra-thin layers. For

Alq3, for example, optimal interlayer thickness is usually less than 10A [1,2,3,4,5,6,7] with

A1 cathodes. Though thicker layers have been used effectively for Ag cathodes [8], device

performance degrades considerably for interlayer thicknesses greater than 30A. For

polymeric conductors with electron conductivity much lower than Alq3, another metal layer,

such as Ca, must be added to the cathode to ensure adequate device performance [9].

-181 -


Chapter 9 LiF interlayer properties on surfaces 182

However, other electron transport layers are able to support much thicker interlayers without

affecting the device performance. Many mechanisms have been proposed to explain the

behaviour o f devices with interlayers o f various thicknesses, ranging from interface doping

by dissociation for thin interlayers [10] to tunneling injection for thick layers [8].

9.1.1 Device behaviour

Previous work within the Lu group [11] has indicated that increasing the LiF interlayer

thickness has a different effect on the luminance characteristics for devices with Alq3 and C6o

electron transport layers. Generally, Alq3 based devices are expected to show poor device

performance with interlayer thicknesses above 40A. As the thickness continues to increase,

the devices typically fail after application o f voltages between 6 and 10V. For an Alq3 ETL,

therefore, the maximum interlayer thickness for device operation is around 60A, as has been

previously reported [8], The case o f C6o is dramatically different from that for Alq3, as C6o

based devices are capable o f operating even with a 100A interlayer [11]. At this thickness,

the tunnelling probability is quite low and LiF should provide an insulating barrier at the

metal/organic interface. This type of behaviour can be related to either the properties o f the

LiF layer itself, such as the thickness or extent o f diffusion of interface, or to the properties

o f the underlying layer, such as the conductivity.

The previously proposed mechanisms for thick interlayers have assumed that this

multicomponent system consists o f sharp heterojunctions, with the interlayer and the

underlying organic as separate layers. Unlike traditional semiconductors, organic/inorganic

interfaces are complex and the interfacial regions can often be thought of as composite

materials with properties very different from each of the individual components. As such, the

conductivity o f the organic layer may also be affected by the presence of such dielectric



materials. X-ray photoelectron spectroscopy is extremely sensitive to changes in surface

conductivity, and can be used as a probe to predict the response of the underlying organic

layer to the deposition of such insulating compounds [12,13,14],

In this chapter, we investigate the behaviour of thick layers o f LiF on Alq3 and C6o

electron transport layers. Unusual charging behaviour observed with X-ray photoelectron

spectroscopy indicates that LiF deposited on Alq3 becomes insulating above nominal

thicknesses of 30A, while LiF on C6o is still conductive even after 100A deposition. This

charging behaviour is less than that observed for LiF crystals, but much greater than that of

thick LiF layers deposited on conducting substrates, such as Si. This observed insulating

behaviour of Alq3-LiF layers can explain a breakdown in device performance above 30A,

whereas C6o-LiF layers, which are still conductive even at extremely high nominal

thicknesses, consistently show superior device performance.

9.2 Experimental

Samples were fabricated on a H-terminated Si (100) wafer, with LiF deposited on either bare

Si or on 1000A of Alq3 or C6o- The LiF thickness ranged from 10A to 100A nominal

thickness, as measured by a quartz crystal microbalance. Films were fabricated by sequential

thermal evaporation in the cluster tool with base pressure o f 1x10' Torr, as described in

Chapter 4. Deposition rates were approximately lA/s, and 0.2A/s for the organics, and LiF

respectively.

X-ray photoelectron spectra were generated using monochromated A1 Ka (1486.7 eV)

radiation and a 23.35eV pass energy. In order to examine the charging behaviour, the beam

intensity was held at 300W for different exposure times. Typical X-ray photoelectron spectra



collection time was ~45min. To examine the transient behaviour of the charging, the samples

were also exposed to the beam for 60min, with spectra collected in 5min intervals. To

compensate for surface charging in the comparison of various samples, spectra were aligned

based on the adventitious overlayer, unless otherwise noted. For FI-terminated Si and LiF

crystals, adventitious C Is core level was aligned to 284.8eV. For the organic layers, where C

Is core level is dominated by signal from the organic molecule rather than adventitious C,

adventitious O Is was aligned to 532.59eV, determined as the average value of adventitious

O after alignment using C Is for LiF on H-terminated Si and LiF crystals. For Alq3, the O Is

was deconvoluted into two peaks, one for Alq3; corresponding to the stoichiometric ratio with

N Is for the Alq3 molecule, and the other assumed to be from adventitious O on the surface.

Comparisons were also made based on spectral alignment from literature values for the

substrate core levels. High-resolution SEM was performed on a Hitachi S4700 at the

Institute for Microstructural Studies at the National Research Council o f Canada.

9.3 Results and discussion

9.3.2 LiF growth on organic surfaces

For an electrical insulator such as LiF, the determination of the absolute binding energy to

establish the chemical state is greatly complicated by charging effects. When charging is

compensated for by alignment to the underlayer, the F Is core levels for the three substrates

show slight scatter in the observed binding energy, as seen in figure 9-1. The core levels were

all aligned to standard literature values for the substrate (for Si [15], for N in Alq3 [16], and

for C6o [17]). For C6o sample, this alignment is complicated by the fact that adventious C

species overlap the only signal from the substrate. For most cases, except for 100A LiF on



C6o, the highest peak in the C Is core level was attributed to C6o- For 100A LiF, the highest

peak observed in the C Is core level was due predominantly to adventitious C, determined by

aligning the O Is core level with that observed on the 50A LiF sample (not shown). From

this alignment with the expected substrate values, though the values for LiF on Alq3 are

~0.7eV lower than that on the other two substrates, all of the F Is values fall within the

acceptable range for ionic LiF [18].

Figure 9-1 F Is core level for LiF of varying thickness on (a) Si (b) C60 and (c) Alq3 aligned to literature values for the underlayer. The solid vertical line at 685.6eV represents the alignment of the core level using adventitious species on the surface for all cases. Due to differential charging effects, the core levels are slightly different for the various substrates, but all fall within the acceptable range for ionic LiF. See text for details. Notice the broad and asymmetric peak shifted to lower binding energies for LiF on Alq3 surfaces attributable to charging effects.

692 690 688 686 684 682 680

Binding Energy (eV)

For insulators, since the Fermi energy of the sample and the spectrometer are

decoupled [19], the overlayer and underlayer may charge differently; therefore, neither the

adventitious species on the surface nor the substrate signal can be used consistently to

reliably retrieve the absolute binding energy for all cases. For LiF on Alq3, for example, the

difference between overlayer and underlayer alignment is 0.4eV; therefore, the appreciable

broadening and shift to lower binding energy values for Alq3 may be attributed to charging as

F 1s 1 0A L iFon Si 30A LiF on Si 50A LiF on Si

100A LiF on Si

(/)»c13.QCD F 1s 0.47

— n— 10A LiF on Ce

— o— 30A LiF on C;.

—a— 50A LiF on Ce

— 7— 100A LiF on C,</)cCD+-•c

F 1 sCDEoz

10A LiF on Alq,

30A LiF on Alq,

50A LiF on Alq,

100A LiF on Alq,

.32eV^



described in section 9.3.4.1. For the Si and C6o substrates, as the thickness increases, the

Fermi level becomes increasingly decoupled and the spread in the F Is values is ~0.6eV.

With 100A LiF on all surfaces, where the influence of the substrate is diminished, the F Is

core levels all align at 685.6eV, using the adventious species on the surface. This is roughly

in the middle o f the range o f binding energy values observed with alignment using the core

levels of the underlayer. Therefore, it appears that LiF is in a similar chemical state

regardless of which substrate it is deposited on.

The only possible exception is for 10A LiF on C6o, which has a binding energy

0.47eV less than that at any other thickness. Neither LiF on Alq3 nor on Si showed any

changes in the F Is core level position upon alignment with the underlayer as thickness

increased. In order to check whether this shift is related to changes in the charge density of

the F Is atoms or final state relaxation effects, the modified Auger parameter may be used as

an excellent tool. As shown in the Wagner map for LiF on these surfaces, figure 9-2, almost

all of the values lie along lines of slope 3, indicating a similar chemical state during growth.

The only exception again is 10A LiF on C6o which shows both initial and final state effects.

>,O)L_CDc

LUoa)ck

■ LiF on Si654.4

• LiF on654.2 a LiF on Alq3

/ i T j — 1654.0

653.8

653.6

653.4

653.2 • a= 13 3 9 .4 3 / / / /s lope= 3

653.0

Figure 9-2 Wagner map for deposited LiF of different thicknesses on Si, C60 and Alq3 surfaces. A majority of the points lie along lines of slope 3 indicating a similar chemical state. The difference between LiF on Alq3 and LiF on the other substrates is likely due to charging effects.

687.0 686.5 686.0 685.5 685.0 684.5

Binding Energy (eV)

The structure of the deposited LiF layer appears to be very similar for any substrate,

with the same thickness o f LiF deposited regardless of the surface. Figure 9-3 shows the



change in the F Is core level intensity with deposition, assuming the same sticking

coefficient on these surfaces (the intensity ratios support this assumption). Assuming an EAL

of 33.5A from a LiF film density of 1.85 g/cm3 [20,21], the predicted layer by layer growth

to 100A would follow the dotted line in figure 9-3. Instead the change in the signal intensity

predicts much thinner films. Though this could be attributed to LiF diffusion into the organic

layers, the growth o f LiF on the molecular surfaces follows a very similar trend to that

observed on Si. Since LiF would be unlikely to diffuse into Si, diffusion into the organics is

also unlikely. Instead the signal intensity behaviour initially suggests that LiF islands are

forming on the surfaces, as has been observed previously [22]. As the change in the F Is

intensity eventually reaches that predicted by a thick layer on the surface, it is likely that the

islands grow in two dimensions with deposition, eventually coalescing into a complete layer.

Figure 9-3 Growth of LiF on surfaces as monitored by XPS. The dotted red line is the expected change in intensity with layer by layer growth, assuming that by lOOA deposition, there is a full 100A LiF layer on the surface. The solid red line represents a linear sum of reduced squares best fit of the data for thicknesses less than lOOA.The intensity follows a parabolic shape (the blackdashed line is just a guide to the eye), indicating initially island growth with eventual formation of a complete layer on all surfaces.

LiF deposition (A)

As figure 9-4 shows, the XPS sputter profile through the depth of the LiF layer is

very similar for both Alq3 and C6o. As can be seen in figure 9-4 (a), the depth at which F is

still visible by sputtering is very similar for 30A LiF on both molecular surfaces. However,

for molecular solids such as C6o and Alq3, which are held together by relatively weak van der

Waals forces, ion sputtering will tend to overestimate the thickness of the LiF layer, as the

sputtered elements are driven down into the soft material. Figure 9-4 (b) shows the sputtering

■ LiF on Si • LiF on

LiF on Alq.

0

■2

■3

■4

■5

0 10 20 30 40 50 60 70 80 1



profile from peeled-off device structures with a lOOA LiF/Al cathode with either molecule as

an ETL. Again, the depth at which F is visible after sputtering suggests a similar thickness of

LiF in both cases, though the C6o ETL appears to have a slightly thinner LiF layer. However,

the sputter depth for 100A LiF on A1 is similar to, and sometimes more than, that for 30A

LiF on the two molecules (figure 9-4 (b) and (a)), suggesting that the organic molecules are

too soft to reliably indicate the LiF thickness by sputtering. However, even for the 100A

LiF/Al surface exposed by peel-off, shadowing effects and calibration o f sputter depth using

a SiC>2 standard limits the accuracy of the sputter derived value o f the LiF thickness

F 1s

Surface side /

F 1s

Surface side

695 690 680685

F 1s 100

80

60Surface side

40

cZJ.Qk_

100F 1s->%80</>

Coc Surface side 60

40

20

695 690 685 680 0 5 10 15 20 25Binding Energy (eV) Binding Energy (eV) Sputtering time (min)

fa) fb) fc)Figure 9-4 XPS Ar+ ion sputtering profiles for LiF with C60 (top row) and Alq3 (bottom row), (a) Profile of F Is core levels with a nominal LiF thickness of 30A on organic surfaces (b) Profile of F Is core levels with nominal thickness of lOOA LiF obtained from Al/LiF cathode surface exposed by peel-off (c) Concentration profile for A1 and F for the structure described in (b).

Nonetheless, the relative thickness of the LiF layer can still be compared for the case

of C6o and Alq3. In sputter depth profiling, the interface between two materials can be taken

as just beyond the 50% cross over point in the relative concentration [23]. From figure 9-

4(c), the two cross-over points are very similar, at 11.8 min for C6o and 12.3 min for Alq3,

which is within the depth resolution error for XPS [23]. The LiF layer thickness values,

therefore, are very similar for the two cases. Using a SiC>2 calibration of 6A/min for the



sputtering depth, the layer thicknesses are approximately 70A and 74A, slightly less than the

nominal 100A as measured by the crystal microbalance.

When the LiF layer is thick enough, other techniques can also be used to confirm the

quartz crystal monitor derived thicknesses. High resolution SEM images o f the cross-section,

shown in figure 9-5, indicate a complete LiF layer on the surfaces of roughly similar

thickness in both cases. To minimize charging effects from bombarding the heavily

insulating layers with charged particles, the Alq3 sample was coated with Pt and the images

were taken at a slight tilt from normal. Subsequently the LiF overlayer thickness of 185A is

overestimated by -5 0 A. As was suggested by the intensity change with thickness, even

though the interface between LiF and the organics is not sharp, minimal diffusion o f the LiF

into the Alq3 or C6o layers is visible.

Although there is significant error in the thickness values, with XPS sputter profiling

underestimating the thickness and SEM overestimating it, the LiF layer thickness o f 100A

does appear to be consistent with the nominal thickness as determined from the crystal

microbalance. It is very difficult to accurately measure the thickness with less deposition;

therefore, by convention, the deposited thickness will be assumed to be the same as the

nominal thickness predicted by the quartz crystal thickness monitor.

Figure 9-5 High-resolution cross-sectional SEM images of lOOA LiF on (a) Alq3 and (b) C60. To accommodate charging, the LiF on Alq3 was coated with Pt, and tilted 4° from normal.



In Figure 9-5 above it can be seen that the interface between the organics and LiF is

not sharp, which could also be influencing the thickness measurements using the various

techniques. Although the interface appears to undulate for the C6o surface especially, the LiF

film itself is actually fairly uniform in thickness for both cases. Figure 9-7 shows the surface

topography as visible in the SEM for LiF on the two surfaces. Especially for C6o, the surface

mimics the morphology of the underlying interface visible in figure 9-6, suggesting that the

LiF layer is actually o f uniform thickness. These surface undulations were not as visible in

the cross-sectional image due to surface effects, which distort the image slightly at the

vacuum/sample interface [24]. It is probable, therefore, that upon peel-off, the LiF surface is

not planar and the topographies o f Alq3 and Cô are different. Therefore, the final predicted

thickness values by XPS sputter depth profile are not exactly the same.

SOOnm

(a) Alq3 (b) C6oFigure 9-6 SEM images of the surface topography for lOOA LiF on (a) Alq3 and (b) CWj surfaces. Samples tilted to 45° to image both LiF surface (top left hand side) and cross-sectional cleavage plane through Si (bottom right hand side).

As the morphologies of the underlying organics are so different, it is possible that the

grain structure o f the LiF overlayer film growing on these surfaces may be different, as the

surface packing ratios would not be the same. Due to the low contrast and charging problems

for LiF, however, the resolution of the SEM was insufficient to resolve any such differences

in the film structure.



As the structure and thickness of the LiF can be taken as basically the same for

growth on these two molecular solids, the relative conductivity o f the underlying electron

transport layer plays a major role in determining the maximum useable thickness for the

dielectric layer. With the considerable injection observed in devices, the combination o f LiF

and Ceo appears to be extremely conductive.

9.3.3 Resistivity effects as observed by XPS

It is difficult to accurately measure the conductivity of thin organic films and molecular

devices using traditional conductance techniques [14, 25]. Generally, the introduction of any

“external” contact to probe the film properties will modify those properties making it

difficult to isolate the electrical behaviour o f the film itself. XPS, due to its high sensitivity to

the conductivity o f the films, can be used as a non-contact method for analyzing the

resistance/capacitance and other electronic properties of thin semiconductor and dielectric

films [26] by examining charging effects on the position o f the core levels (surface charge

spectroscopy [27] or chemically resolved electrical measurements [13,14]). Although the

parameters of such a technique are not fully established, and values derived from such

analysis may not be the absolute conductivity values, an examination o f charging effects

using basic electrostatic theory can give insight into the relative electrical properties and help

to explain the observed device behaviour of Alq3 and C6o with thick LiF films.

9.3.3.1 Charging effects in XPS

Charging effects result from the competition between electron emission into the vacuum and

electron redistribution from the surroundings that imperfectly compensate for the emission

(see Chapter 3). In insulators, charge build-up at the sample surface retards outgoing



photoelectrons and increases the apparent binding energy. For this study, since the change in

surface potential during irradiation is o f interest, the “charging shift” is defined by alignment

to the adventitious overlayer. Unlike the case o f oxide growth on a metal, where differential

charging between the substrate and the overlayer may be observed on the same peaks [26],

this differential effect cannot be used to compare the observed charging shifts for different

substrates. Instead, the absolute charging is used as an indication o f the surface potential.

9.3.4 X-ray induced charging effects fo r LiF coatedfilms

As an insulating material, LiF is expected to show some charging effects, especially with a

monochromated X-ray source [19]. This charging is expected to increase with thickness, as

shown in figure 9-7 [26,28].

Figure 9-7 Observed charging shift as a function of thickness.Lines represent linear sum of reduced squares best fits of the data. Assuming a parallel plate model, the slope of the lines represent the electric field developed in the dielectric, given on the graph in units of MV/cm.

0 10 20 30 40 50 60 70 80 90 100 110


As a first approximation, the relative induced charge in the sample can be estimated

from the shift relative to the thickness from basic electrostatic theory, assuming a parallel

plate capacitor (equation 3-10). Note that this may not be the real value, as the charging shift

appears to become independent o f thickness for LiF on Si surfaces at thicknesses greater than

50A (both 100 and 200A LiF on Si shift by the same amount in figure 9-7). Here Vs, the

observed potential at the surface can be represented by the charging shift as:

• LiF on Si » LiF on

r LiF on Alq,3.5

3.0

> 2.5a)c 20

3.13MV/cm

200A on Si

1.86MV/CI

0.51.65MV/cmro o.o.co

-0.5



A Ey _ chKgmg ê

where e is the fundamental electron charge.

As can be seen from figure 9-7, the slope o f the charging behaviour observed for LiF

on Alq3 surfaces is about twice that of the other two surfaces. As a first approximation,

assuming that the dielectric constant is the same for LiF in all three cases, the induced charge

within the LiF layer on Alq3 is about twice that observed on C6o or Si.

Differences in surface charging could be attributed to the formation of much thicker

films or higher surface roughness [29]. However, as described above in section 9.3.2, the LiF

thickness is the same for all substrates. Additionally, due to the nature of the underlying

organic films, a LiF layer on Alq3 is much smoother than one of comparable thickness on

C6o, as observed by high resolution SEM (figure 9-6). The RMS roughness (as determined by

AFM measurements) for LiF on C6o is 10 times larger than on Alq3. Consequently, neither of

these can explain the high charging on LiF/Alq3 surfaces.

Potentially, charge compensation mechanisms [28,30] other than the bulk film

conductivity, as described in section 3.4.1, may also be different for Si and C6o compared to

Alq3. For the case of LiF on the organic surface, many of the possible compensating

processes, however, are actually the same. Compensating electrons from the surroundings

would be virtually unchanged for all experiments as the same acquisition gun and sample

holder were used in the same chamber with nearly the same pressure; therefore, any electron

or secondary electron emission due to [28] the gun material, the chamber walls, the X-ray

window (not applicable for monochromator use), the gauges and pumps, the vacuum pressure

and composition would also be the same for all cases. The irradiation conditions - the

incident photon flux, the irradiated area and photon energy - are nearly constant for all



experiments, and can be eliminated as giving rise to potential differences. Variations in

photon flux were examined specifically by leaving the source running continuously during

acquisition from different samples, and differing charging behaviour was again observed. It

is also unlikely that the observed differences could be due to poor contact between the

sample and the sample holder, since all the samples were deposited on the same Si wafer

substrate and mounted with the same conductive C mounting tape. Other factors [30] that

may affect the charging behaviour that can be taken as similar for these systems are the

photoelectron cross section, the effect o f electron emission from the solid on the effective

attenuation length, the electron yield (electron emission distribution per photon), and the

effective retention coefficient/electron affinity o f surface, since the same core level was

examined from samples o f the same thickness, lateral dimensions, surface cleanliness,

temperature, and nominal composition.

Having eliminated all the possible external and internal factors that could account for

charge compensation within LiF deposited on Si or C6o, the only other explanations for the

unusual charging behaviour are different induced photocurrents or conductivities for the

underlying organic layer, assuming constant bulk conductivity o f the LiF layer. The impact

o f electron excitation in the underlayer is also supported by the fact that LiF/C6o shows

negative charging shifts with thin LiF layers, as seen in figure 9-7. When the LiF thickness

is less than 30A, more photoexcited electrons from the C6o layer are able to accumulate at the

LiF surface than are lost from the LiF layer itself [31], and the charging shift is negative.

If the surface charge density can be attributed mainly to trapped holes, the resultant

density o f hole traps can also be calculated, as in table 9-1, in section 9.3.4.2. Charging is not

a static or instantaneous process, as it is related to the neutralization of the built-up charge by

charge carriers from the substrate material. A crude estimate o f the conductivity of the bi



layer structure under irradiation can be made by examining the core level positions with

irradiation time if there are no chemical changes due to X-ray bombardment.

9.3.4.1 Transient effects

As figure 9-1 of section 9.3.2 above shows, the F Is core levels for LiF on Alq3 are

broadened and asymmetric compared to those o f LiF on the other substrates. Also, the F Is

core level for LiF on Alq3 shifts visibly within the time scale of the experiment. To examine

this transient behaviour, the samples were examined in 5min intervals under a constant 300W

X-ray beam for 60min.

Generally, broadening and peak shifting under irradiation is an indication that the

material is unstable under X-ray bombardment. Alq3 itself is a fairly conductive molecule,

and shows no tendency to degrade chemically under X-ray irradiation during the time scale

o f the experiments, as has also been reported by other groups [32,33], There were also no

apparent charging effects, seen by a shift in the A1 2p core level, within the time scale of

these experiments (see figure 9-10), and have never been observed previously, not even for

thick layers on insulating glass substrates [34], This does not imply that there is no charge

build-up in Alq3 itself, only that the steady state value is reached on a time scale less than the

acquisition time [26]. X-ray bombardment has been shown to affect LiF [35], through the

generation o f F-centres. However, it is unlikely that the observed shift can be attributed to

this degradation mechanism, which tends to decrease the observed charging with time, rather

than increase it.

To confirm that the observed behaviour can be related solely to charging effects, the

difference in the Li-F Is binding energy and the composition ratio were examined as a

function of irradiation time, as shown in figure 9-8(a) and (b). Though the crystal and



deposited layers do show different A(F 7s-Li Is), none are significantly affected by

irradiation. It is unclear at this time why the values are so different. One possible explanation

is that the polycrystalline samples are highly disordered and there is a difference in the

Madelung potential difference which affects the binding energy positions. Another

possibility is the impact o f the static charging shift, which is ~50eV for the crystal. Due to

the broad and asymmetric nature of the peaks (FWHM ~3.5eV), the error in the binding

energy position for each element may be greater than leV. This is not quite sufficient to

encompass the difference between the crystal values and those of the deposited layers.

Further investigation o f this effect may be required. Nonetheless, the chemical state o f the

crystal, and the deposited layers, though different are not significantly affected by the

irradiation.

The composition (figure 9-8 (b)) does decay slightly for the crystal, indicative o f the

formation of defect F-centres [35]. Photon stimulated desorption is well known to change the

stoichiometry o f fluorides [36,37]. As the deposited LiF layers are likely polycrystalline, the

defect structure formation does not have a visible effect even though there is a slight decrease

in the composition.

633.0

632.5

632.0

631.5

- □- 200A LiF on Si .- a - 50A LiF on Alq3

- v - 100A LiF on CM .

- o - LiF crystal

631.0

111 630.5 CQ< | 630.0

629.5 -1 zw-629.0

0 10 20 30 40 50 60

03

1.5

1.4

1.3

1.2

1.1

1.0

0.9

0.8

0.7

I 1 I ' I- □- 200A LiF on Si- a - 50A LiF on Alq3

- v - 100A LiF on Cw- o - LiF crystal

T )

1 1

10

Time (min)20 30 40

Time (min)50 60

Figure 9-8 (a) The AE F ;5.Li ]s over time indicating that the chemical state is consistent with irradiation time, though LiF crystal is different from the deposited layers, (b) Change in the Li/F ratio over irradiation time indicating a slight decay due to the formation of F-centres The lines are just a guide to the eye.



This suggests that the chemical environment is not changing during irradiation, and

that the observed shifting in the binding energy can be attributed to charging effects. A

similar argument can be made for LiF on C6o and on Si, though they do not show appreciable

broadening. Since the observed effects can be attributed to the conductivity of the layer, the

change in core level positions with time can be used as a crude estimate o f the conductivity

o f the film with respect to charge trapping processes.

The surface potential change during irradiation, however, depends on photoelectron

emission and various relaxation processes in addition to the time dependent transport of

neutralizing charge carriers [26], Many o f these processes occur on much faster time scales

than that of XPS measurement, so the observed potential at the surface consists o f two parts,

^ c h a r g i n g = ^ initial + A£’(/) (9-2)

As only the transient portion is of interest here, the transient charging shift is defined as the

difference from the F Is position at t=0 o f acquisition.

9.3.4.2 Estimation offilm resistivity from transient effects

The accumulation of charge within the insulating layer is related to the conductivity o f the

underlying substrate layer; therefore, the evolution of the charge distribution with time is

related to the thickness o f the overlayer. Above a critical thickness, the X-ray stimulated

electrons in the underlayer are no longer able to tunnel through the LiF layer to prevent

significant charging at the surface. As the spectrometer resolution was around 0.33eV (see

Appendix C), the core level shifts can be considered resolvable if they are greater than

~0.3eV. As can be seen in figure 9-9, for very thin layers regardless o f the substrate, there are

negligible or even negative shifts in the core level position over time, which can be



neglected. At a certain thickness, however, the observed charging suddenly becomes

thickness dependent. This thickness dependence is very different for the three substrates.

Figure 9-9 Shift in the F Is after 45min X-ray bombardment for LiF on various substrates. The lines represent linear sum of reduced squares best fits of the data above a critical charging shift. The cross-over point is indicated for each curve.

" " o 10 20 30 40 50 60 70 80 90 10 (T l90200210


If the critical thickness for charging is defined as that at which a 0.5eV shift is

observed, then an empirical relationship between the conductivity of the underlying layer and

the critical thickness can be established. Interestingly, it appears that the hole mobility gives

a fairly good indication of the critical thickness. This is not unreasonable, as the resultant

majority charge carrier with X-ray bombardment are holes, left behind after the

photoelectrons are excited into the vacuum. As the conductivity is a combination of the

mobility and the number o f charge carriers, this indicates that the critical thickness can be

determined by the conductivity o f the underlying substrate layer, since the number of charge

carriers induced by X-ray excitation should roughly be the same.

The X-ray irradiation induced charging observed in XPS is in some ways analogous

to the expected electric charge build-up during device operation [26]. Therefore, the observed

charging behaviour may be able to explain the extreme difference in the useable range of LiF

thicknesses for Alq3 and C60 based devices. From the work of Huang [11], Alq3 based

devices with a 60A LiF layer appear to fail after a minimal stress of 8V, well below the

operating range for a typical display. If one assumes that the corresponding charging shift

* LiF on Cgo LiF on Alq.> 2.0

.5

.0

61.1.5

21.0 127.1

0.0



visible with XPS analysis would represent the upper limit for adequate charge movement

through the coupled LiF/organic layer, then a similar charging shift for C6o would require

~190A of LiF.

The charging response of the different materials can be related to a number of factors

such as the secondary electron yield of the underlying material1, the LiF film diffusion into

the organic layer, and the LiF film integrity. While the secondary yield of Si would be

approximately twice that o f the organic films, and could be used to explain the different

charging effects, the two organic films have very similar phonon and electron eliminating

mean free paths, giving roughly the same relative yield as estimated using the analytical

equation of Henke et al. [38]. As described in section 9.3.2, SEM and XPS indicate that very

minimal diffusion is occurring in either o f these systems. Therefore, one possible explanation

for the observed differences could be differing grain structures o f the LiF film on the two

different surfaces. The limits of the resolution for SEM for such light elements make it

difficult to examine the LiF grain structure directly.

In order to compare the electrical properties of the LiF/underlayer combination, care

has to be taken to choose a LiF thickness that should provide a similar charging response for

various substrates. From figure 9-9 above, 50, 100 and 200A LiF on Alq3, C6o and Si

respectively, are appropriate systems for comparison.

As seen in figure 9-10, thick (200A) layers of LiF on conductive Si reach a steady-

state condition fairly quickly, representing normal charging behaviour for an insulator atop a

conductive material. The irradiation induced charging shift from even 50A LiF on Alq3 is

higher than that for LiF on Si, but not as extreme as that observed for insulating LiF crystals.

1 Here, secondary electron yield refers to all the electron excited by X-ray bombardment including photoelectrons, Auger electrons and electrons that have lost some energy through scattering.



On C6o, a 100A of LiF are needed to produce shifts larger than those for LiF on Si, but this

LiF/Ceo combination still does not show as much charging as 50ALiF on Alq3.

As a first approximation, the transient relationship may be described by an

exponential decay using equations 3-15 and 9-3 as [26, 39],

where AE(°°) is the steady state value of the binding energy shift taken at t=60min, based on the secondary electron flux from the material, and r is a thickness dependent time constant.

Though the secondary electron flux also changes with charging, equation 9-3 is

adequate as a first approximation, and is consistent with the change in the F Is position. The

of the data with the model function (equation 9-3) with Levenberg-Marquardt statistics.

Table 9-1 lists the calculated time constants; the developed electric field, determined from

the slope of the best-fit lines in figure 9-7; the estimated charge density, assuming a parallel

plate capacitor with a relative permittivity, 8LiF, of 9 .0368o for LiF [40]; and estimated

conductivity of the LiF layer on the different substrates and for the LiF crystal.

(9-3)

time constants and steady state binding energy shift can be estimated from a reduced chi fit

Figure 9-10 Shift in F Is core level with irradiation time. The lines represent reduced chi2 fits of the data to a function described by equation 9-3 with Levenberg-Marquardt statistics. The right facing triangles for Alq3 indicates the change in the N Is core level with time to indicate the stability and conductive nature of the molecule itself.

0 10 20 30 40 50 60

Tim e (min)



Table 9-1 Estimated electrical properties of the LiF film and crystal (first approximation -equation 9-3)

Structure F.(\j Yc m! Q (/cm )*x!0 ^ s x 10 T ( s ) A E ( qo) 5^(,|L.nlL.nso )xl0

LiF/Si 1.86 1.49 9.30 353 0.67262 22.7

LiF/C60 1.65 1.32 8.25 498 0.95138 16.1

LiF/Alq3 3.13 2.5 15.7 1447 1.69057 5.53

LiF crystal — — — 790 3.01233 10.1

* £ 'l 1f = 9 .036 £ o § Sc ^

If the charging behaviour were in fact to follow an exponential decay as has been

assumed, then the charge lifetime as determined above should be the same over the whole

time scale. From figure 9-11, there are clearly at least two, and possibly three, regions with

different time constants. The values derived from the non-linear curve-fitting only fit the

initial portion o f the data. Therefore, the simple exponential decay model is insufficient to

describe the charging behaviour, due to the added complication that the charge lifetime is not

constant over the whole time scale.

Figure 9-11 Change of the F Is core level kinetic energy as a function of time. For each set of data, the first set of lines represents the time constant derived from the non-linear curve fitting to figure 9-10. The other lines represent linear sum of reduced squares best fits of the data for the various regions.

0 10 20 30 40 50 60

Time (min)

This suggests that there may be multiple kinds o f traps for holes within the

LiF/underlayer system [26]. Initially, the material dependent hole traps dominate the

charging. Once these are all filled, traps with longer lifetimes become prominent in the

a 200A LiF on Si 50A LiF on Alq

^ 100A LiF on C6

o LiF crystal

1

0

1

S -2

■3

■4



charging behaviour. The lifetime o f this second type of traps is very similar for every case,

even the LiF crystal, indicating that it is due to charge trapping within the LiF layer itself.

The behaviour o f LiF on Alq3 is dominated by these long lifetime traps from the outset, with

a visible change in the time constant after 40 mins, similar to the behaviour in the crystal.

Generally, a change to shorter filled traps is consistent with ordering over time. This

behaviour implies that over time, due to the built up electrical field, the LiF dipoles may

become oriented and allow for faster charge movement. Contrary to this, a much thicker

layer o f LiF on C6o initially has much shorter lived charges, with a visible change in the time

constant after 5 mins of irradiation, similar to Si in behaviour. One possible mechanism for

this type of behaviour is relaxation o f the LiF dimers related to the differing morphologies on

the LiF/C6o and LiF/Si surface. As figure 9-11 shows, the behaviour o f LiF on C6o layer is

very similar to that on a conductive surface, whereas that of LiF on Alq3 is very similar to the

insulating crystal.

In order to compare the relative conductivity, it is useful to determine an effective

time constant over the whole time scale. As the secondary electron yield is a function o f the

surface voltage, not a constant as assumed in the first approximation, a much better estimate

o f the effective time constant can be made from the initial and steady-state solutions to the

change of surface potential with time (equation 3-14).

In the initial state, the surface potential, and hence the shift in the F Is binding energy

is approximately linear with time,

3AEf „ F S eS ’ j 0) dt C

where AEF /.v is the shift in the F Is core level, S is the uniformly irradiated specimen surface area, <P is the X-ray irradiation flux, S x(0) is the initial secondary electron yield, and C is the geometric “capacitance” defined by the ratio between the F Is binding energy shift and Q+S.



C can be estimated from the linear portion o f the transient shift in the F Is core level.

At the steady state (t = o ° ) ?

AE'--~ — = F S eS x{Vs ) (9-5)R

where AEf /ft00) is the steady state shift in the F Is core level, R is a resistance value incorporating both sample resistivity and any self-regulation effects, and 5 ftVs) is the secondary electron yield at t=°°.

Therefore, R can be estimated from the F Is transient shift curve, assuming that the system

has reached the steady state potential by the end o f the experiment. Though this assumption

may not be strictly true for LiF on Alq3 , the value o f the surface potential at 60 mins was

taken as the steady-state value in all cases. These estimations are shown in figure 9-12.

Using these equations, and assuming an equivalent circuit description o f the charging

as described in section 3.4, the effective time constant, t, can be defined as the product o f the

resistance, including both sample resistivity and all the self-regulation effects coming from

the vacuum [28], and the “capacitance” defined by the ratio between Vs and Q+S, i.e. x=RC.

Figure 9-12 Estimation of the R and C values from the linear and steady state portions of the transient F Is core level shifts. Lines are just a guide to the eye.

0 10 20 30 40 50 60

Time (min)

As the X-ray flux, irradiated area and initial electron yield from LiF are the same for

all cases, the time constants give an indication of charge lifetime within equivalent systems.

For this system, the secondary electron yield from LiF can be taken from Henke et al. [39],

£ . . \J 1 ! . 1-—o—200A LiF on Si — a — 50A LiF on Alq3 - v - 100A LiF on CBn

“ l ■-------- 1-------- 1-------- 1“

V H1.5

>CD

1.0 - A V /A t ,

-Cww

0.5 -

0.0 -

V ( o o ) -

□—□—10- -□—a- n— 0



scaled to the present irradiated area, or estimated from the analytical formula developed in

the same work2. Table 9-2 below lists the second approximation values for the electrical and

dielectric properties o f the thin film systems.

Table 9-2 Estimated conductivities for LiF thin films and crystal from the transient F Is core level shift (2nd approximation)

Structure r -12 L (F )x10 p 14* tM£2)xl0 T (s)

„ -1 -15§ S (Q )x l0

200A LiF/Si 1.46 4.67 681 53.5

100ALiF/C60 1.24 7.15 888 17.5

50A LiF/Alq3 1.60 13.4 2139 4.67

LiF crystal 0.46 39.3 1820 44.0* — P C § <j = T , assuming a parallel plate capacitor [ Ch

X ~ K L r { ~ S

As can be seen in the above table, the derived capacitances for the films o f different

thicknesses on the various substrates are very similar, and much higher than that o f the LiF

crystal. Therefore, the original assumption that the dielectric constant was the same for all

three cases and equal to that o f a LiF crystal was incorrect. The value of the effective

dielectric constant can now be determined from the experimental capacitance value. Since

the penetration depth for X-rays is much larger than the LiF thickness in all cases, as was

predicted by the thickness dependence of the charging shift, the interaction between the LiF

overlayer and the substrate material is also controlling the dielectric properties o f the system.

Assuming that the LiF/substrate combination can be represented by a series capacitor model

[41], the effective dielectric constants can be determined by

d _ hLiF | d - hUF

^ e f f ^ L i F ^ su b s tra te

where d is the penetration depth of X-rays (2000A), hup is the deposited LiF thickness and £; is the static dielectric constant for material i.

2 See Appendix D.2 for details.



Though the predicted dielectric constants also underestimate this effective value for

all cases, the normalized value of the predicted and calculated dielectric constants with that

for LiF/Alq3 does suggest that the system may be represented by a series capacitor model,

and the calculated values o f the dielectric constant is due to the combination o f the properties

o f the LiF and the underlying layer. The charge density built up within the total LiF/organic

layer can therefore now be calculated using the new value of the dielectric constant. All o f

these dielectric properties are summarized in table 9-3.

Table 9-3 Dielectric properties for LiF thin films on various substrates

Structure effective normalized £ exp normalized n + 2 -5 L? (/cm )xl0 ^ x l O 14

200A LiF/Si 11.11 3.64 4.12 3.65 6.78 4.24

100A LiF/C60 4.52 1.48 1.75 1.56 2.56 1.60

50A LiF/Alq3 3.05 1.00 1.13 1.00 3.12 1.95

From this analysis, it appears as if there are approximately the same number o f hole

traps per unit area for LiF deposited on the two organic surfaces, as would be expected. Since

the capacitance of the layers were similar, and the number of hole traps within the LiF layer

are similar, the observed charging behaviour is likely due to the change in the resistance of

the combination. Since the LiF layer is the same for all substrate, this difference in the

resistance is solely a function of the conductivity o f the underlying layer, as seen in figure 9-

13, where the observed resistances are given as a function of the electron mobility o f the

underlying substrate.



y=7.73133-1 .0227e '14x

CD -(

0.4 0.6 0.8 1.0 1.2 1.4

Figure 9-13 The calculated resistance as a function of the charge carrier mobilities. As the capacitance of the systems are very similar, the conductivity is related to changes in the resistance of the underlayer. The line represents a linear sum of reduced squares best fit of the electron mobility data.

R esistance x1015 (£2)

T h e t im e c o n sta n ts ca n th e n b e u s e d as a n in d ic a t io n o f c h a r g e m o b il i ty in th e

m ateria l, e v e n th o u g h th e o b se r v e d t im e c o n sta n t i s n o t a tru e m e a s u r e o f th e d ie le c tr ic

r e la x a tio n [26]. A s ta b le 9-2 in d ic a te s , th is a n a ly s is ca n s t ill a l lo w s o m e s ig n if ic a n t

c o n c lu s io n s to b e d ra w n a b o u t th e d iffe r in g d e v ic e b e h a v io u r o b s e r v e d fo r C (,o an d A lq 3

b a se d d e v ic e s w ith th ic k in ter la y ers . W ith 100A L iF o n C6o, th e su r fa c e s h o w s a lm o s t fo u r

t im e s th e c o n d u c tiv ity o f 50A L iF o n A lq 3 . T h er e i s a s im ila r im p r o v e m e n t in th e d e v ic e

p ro p ertie s , as ca n b e s e e n in f ig u r e 9-14, w h e r e th e d r iv in g v o lt a g e fo r a d e v ic e w ith 100A

L iF o n Cgo is m u c h le s s th an that o f 40A L iF o n A lq 3.

10000

8000N

E" o 6000

<D OC 4000 CO cE 2000

_l0

0 2 4 6 8 10

Bias (V)Figure 9-14 Comparison of device behaviour for 40 and 100A LiF interlayers with C60 and Alq3 based devices (adapted from [11])

— 40ALiF/C60

— o— 1OOALiF/C60

— a — 40A LiF/Alq.

—a— 100A LiF/Alq.


Chapter 9 LiF interlayer properties on surfaces

9.4 Summary

207

The LiF layer itself is the same regardless o f the substrate on which it is deposited. The

observed differences in the behaviour both in XPS analysis and in the device can be

attributed to the nature o f the underlying film. The C6o layer can be considered as behaving

as if it were a metal. The Al/LiF/C6o could now be thought of as a Metal-Inorganic-Metal

capacitor, and electron injection occurs at the “floating” electrode (i.e. C6o) and Alq3

interface. As long as the LiF layer is thin enough to allow adequate conduction o f free

electrons, the LiF layer can act as a charged source o f electrons to be injected into the

emission layer, in a behaviour similar to a flash memory device. This is also very similar to

the behaviour already observed with metal-organic-metal (MOM) cathodes [42], where the

electron transporting organic layer can be thought of as the insulator. The maximum useful

thickness o f the LiF layer can be directly related to the conductivity o f the bottom contact.

With thin layers o f LiF, Alq3 is also sufficiently conducting to act as its own floating

electrode for injection into the emission layer. The maximum usable thickness for the LiF

would be that at which the potential drop across the capacitor was too great to ensure

adequate injection, and the device would fail. C6o, with much greater conductivity than Alq3,

is able to accommodate a greater potential drop, and a much thicker LiF layer is still

effective. This type of description of the cathode can also help to explain some o f the

behaviours observed by other investigators with different multilayer cathode configurations.

If the underlayer is replaced by an even more conductive material, such as a metal, organic

films with lower conductivity can still benefit from having a LiF layer at the interface [9].

Similarly, if the LiF is replaced with another insulator, the device often performs similarly

but the optimal interlayer thickness changes [43]



9.5 References

'M. Matsumura, K. Furukawa, and Y. Jinde, Thin Solid Films 331, 96 (1998).

2 L. S. Hung, C. W. Tang, and M. G. Mason, Appl. Phys. Lett. 70, 152 (1997).

3C. Ganzorig, K. Suga, and M. Fujihira Mater. Sci. Eng B 85,140 (2001).


5M. G. Mason, C. W. Tang, L-S. Hung, P. Raychaudhuri, J. Madathil, D. J. Giesen, L. Yan,Q. T. Le, Y. Gao, S-T. Lee, L. S. Liao, L. F. Cheng, W. R. Salaneck, D. A. dos Santos, and J. L. Bredas, J. Appl. Phys. 89, 2756 (2001).

6T. Wakimoto, Y. Fukuda, K. Nagayama, A. Yokoi, H. Nakada, and M. Tsuchida, IEEE Trans. Electron Devices 44, 1245 (1997).

7M.B. Huang, K. McDonald, J.C. Keay, Y.Q. Wang, S.J. Rosenthal, R.A. Weller, and L.C. Feldman, Appl. Phys. Lett. 73,2914 (1998).

8 X. J. Wang, J. M. Zhao, Y. C. Zhou, X. Z. Wang, S. T. Zhang, Y. Q. Zhan, Z. Xu, H. J.Ding, G. Y. Zhong, H. Z. Shi, Z. H. Xiong, Y. Liu, Z. J. Wang, E. G. Obbard, X. M. Ding,W. Huang, X. Y. Hou, J. Appl. Phys. 95, 3828 (2004).

9 T. M. Brown, R. H. Friend, I. S. Millard, D. J. Lacey, J. H. Burroughes, and F. Cacialli,Appl. Phys. Lett. 79,174 (2001).

10 M.G. Mason, C.W. Tang, L.-S. Hung, P. Raychaudhuri, J. Madathil, D.J. Giesen, L. Yan,Q.T. Le, Y. Gao, S.-T. Lee, L.S. Liao, L.F. Cheng, W.R. Salaneck, D.A. dos Santos, and J.L. Bredas, J. Appl. Phys. 89, 2756 (2001).

11 C.J. Huang Internal Report LG403 (2004).

12 G. Ertas, U. K. Demirok, A. Atalar, and S. Suzer, Appl. Phys. Lett. 85 183110 (2005)

13 H. Cohen MRS Fall Meeting 2005 Symposium I Interfaces in Organic and Molecular Electronics 17.1; (b.) M. Dubey, I. Gouzman, S. L. Bemasek, and J. Schwartz, MRS Fall Meeting 2005 Symposium I Interfaces in Organic and Molecular Electronics 17.2

14 H. Cohen Appl. Phys. Lett. 85 1271 (2004).


16 A. Turak, D. Grozea, X.D. Feng, Z.H. Lu, H. Aziz, A.-M. Hor, Appl. Phys. Lett. 81, 766 (2002).1 7 J.A. Leiro, M.H. Heinonen, T. Laiho, and I.G. Batirev. J. Elec. Spec, and Rel. Phen, 128

205 (2003).18 NIST X-ray Photoelectron Spectroscopy Database - Version 3.4 (Web version) , National

Institute of Standards and Technology, Gaithersburg, MD, (2003).

19 R.T. Lewis and M.A. Kelly, J. Elect. Spect. Rel. Phenom. 20 105 (1980).20 D. B. Sirdeshmukh, L. Sirdeshmukh, and K. G. Subhadra, Alkali Halides: A Handbook o f

Physical Properties (Springer series in Materials Science) (Springer-Verlag, Berlin, 2001),Vol. 49, p .l.



21 M. Ohring, The Materials Science o f Thin Films (Academic, Toronto, 1992), p. 232-233.

22 T. Yokoyama, D. Yoshimura, E. Ito, H. Ishii, Y. Ouchi, and K. Seki, Jpn. J. Appl. Phys., Part 1 42, 3666 (2003).

23 S. Hofmann, in Practical Surface Analysis, 2nd edition, edited by D. Briggs and M. P. Seah (John Wiley, New York, 1990), Vol. 1, Chapt. 4, p.143-199.

24 J. Fraser, Private communication.

25 Z. J. Donhauser, B. A. Mantooth, K. F. Kelly, L. A. Bumm, J. D. Monnell, J. J. Stapleton, D. W. Price, Jr., A. M. Rawlett, D. L. Allara, J. M. Tour, and P. S. Weiss, Science 292, 2303 (2001).

26 S. Iwata and A. Ishizaka, J. Appl. Phys. 79 6653 (1996).

27 W. M. Lau Appl. Phys. Lett. 54 338 (1988).



30 C.D.Wagner J. Elect. Spect. Rel. Phenom. 18 345 (1980).

31 G. Beamson and D. Briggs, Surf. Inter. Anal. 26, 343 (1998).

32 L. S. Liao, L. S. Hung, W. C. Chan, X. M. Ding, T. K. Sham, I. Bello, C. S. Lee, and S. T. Lee, Appl. Phys. Lett. 75, 1619 (1999).

33 Y. Park, V.-E. Choong, B. R. Hsieh, C. W. Tang, T. Wehrmeister, K Mullen, and Y. Gao, J. Vac. Sci. Tech. A. 15 2574 (1997).


35 G. Johansson, A. Hedman, A. Bemdtsson, M. Klasson, and R. Nilsson, J. Elect. Spect.Rel. Phenom. 2 295 (1973).

36 W. Eberhardt, Phys. Rev. B 46, 12388 (1992).IT

S. Tanaka, M. Mase, M. Nagasono, and M. Kamada, J. Electron Spectrosc. Relat. Phenom. 92, 119(1998).

38 B. L. Henke, J. Liesegang, and S. D. Smith, Phys. Rev. B 19, 3001 (1979).

39 J. Cazaux, J. Appl. Phys. 59, 1418 (1986).

40 C. Andeen, J. Fontanella, D. Schuele Phys. Rev. B 2 5068 (1970)

41 B. Chen, H. Yang, L. Zhao, J. Miao, B. Xu, X. G. Qiu, B. R. Zhao, X. Y. Qi, and X. F. Duan, Appl. Phys. Lett. 84, 583 (2004).

42 X. D. Feng, R. Khangura, and Z. H. Lu, Appl. Phys. Lett. 85, 497 (2004).

43 B. DAndrae, H. Yamamoto, M. Rothman, M.-H. Lu, and J. Brown, MRS Fall Meeting Symposium I Interfaces in Organic and Molecular Electronics, Boston, (2005), II 1.6.


Chapter 10

Interfacial structure models and conclusions

10.1 Introduction

The results outlined in chapters 5-9 indicate that the interface formation process is complex.

To obtain a complete, accurate picture of the interfacial structure, the combination of two

methods is extremely useful: 1. the unique technique o f peeling apart fabricated devices

under high vacuum to analyse both sides o f the buried interface and 2. the more traditional

approach of growing and analysing monolayers of one material grown atop another. By using

this combined approach, one may make connections between the interfacial structures in

manufactured devices and those observed during traditional surface science investigations.

Ultimately this allows a connection to be made between the cathode/organic interface

structure and the behaviour o f devices.

- 2 1 0 -


Chapter 10 Interfacial structure models and conclusions 211

In this chapter, the major findings from this dissertation are summarized and

schematic models of the interface structure between the various cathode materials and active

organics are established. One o f the critical insights gained using the combined approach

outlined in this dissertation is that a universal cathode is likely not possible for every organic

semiconductor. The role of the interfacial dielectric layer, which may even exist at simple

metal/organic interfaces due to interfacial reactions, changes depending on the interactions

that are possible between the differing materials.

The cathode/organic interface generally cannot be considered as a simple junction o f

two (or more, if multilayer cathodes are used) materials. The deposited metal layer,

interfacial layers, and the organic underlayer must all be considered to fully describe the

interfacial structure. Therefore, when organic molecules are modified to improve device

performance, a standard cathode may no longer be suitable. Rather than assuming a universal

cathode, certain criteria for the interfacial structure and stability must be investigated for

every new combination of materials to select suitable cathode candidates. As this dissertation

has outlined, some of these criteria can be determined from simple material property

information, such as the bulk lattice constants or conductivity, or by assuming inorganic

analogues for organic molecules.

The major conclusions of this investigation, described in greater detail in the

following sections, are that

• the interfacial reaction chemistry for an organometallic such as Alq3 may be

predicted by assuming AI2 O3 as an inorganic analogue. Using this analogue,

molecular fragmentation may be described as a metal-exchange oxidation-

reduction type reaction.

• at the thicknesses, 5-10A, typically used for OLEDs, LiF has a metal-dependent

impact on the oxidation behaviour of metals, protecting A1 from oxidation, but

accelerating the formation of carbonates for Mg.



• lattice matching between the dielectric and the metal may be used as a guide to

oxidation behaviour even for amorphous/polycrystalline systems.

• the oxidation behaviour o f a bi-layer cathode incorporating LiF in an OLED is

controlled by the LiF-metal interaction.

• LiF tends to form a charge transfer complex with electron transporting organic

molecules.

• LiF does not appear to follow a layer-by-layer growth mode, irrespective o f the

substrate.

• the maximum useable interlayer thickness for devices can be predicted using the

charging behaviour observed by XPS analysis.

• bi-layer cathodes such as Al/LiF on organics should really be considered as metal-

inorganic-“metal” capacitors, with the maximum useable LiF thickness related to

the conductivity o f the organic electron transport layer.

10.2 Metal/Alq3 interfaces

The interface formation may be described quite easily for relatively simple combinations of

metals and an organometallic such as Alq3 . Figure 10-1 shows a schematic summary o f the

observed interfaces in metal/Alq3 systems.

Mg.Ag

MgO*,Mg, Ag, Al

Ag

M g O x , ; A [ : : : : : : :

Ma,:Ahjs

Au

Figure 10-1 Schematic of various interface structures



Ag and Au cathodes do not react with Alq3 . The interface, therefore, can be described

as a physisorbed (non-reactive) diffuse interface with layers of cathode and Alq3 on either

side o f the diffusion layer (figure 10-1(c) and 10-1(d)). For Mg based cathodes, interface

formation with Alq3 follows a rather complex reaction/diffusion process (figure 10-1(a) and

10-1(b)).

For Mg/Alq3 , the interface consists of a diffusion layer into the organic side, and a

reaction/diffusion layer into the cathode side o f the junction. For the Mg: Ag/Alq3 interface,

the junction has only a single reaction layer. In both cases, the buildup o f metallic Mg at the

interface suggests that the oxidation o f Mg and reduction o f A1 in Alq3 is limited by the Mg-

Alq3 reaction rate. For Mg:Ag alloy cathodes, the presence o f both metallic and oxide species

at the interface indicates that the interface is not as sharp as for pure Mg. The slightly limited

Mg diffusion due to the presence o f Ag has little effect on oxide formation, since the Alq3

fragmentation reaction is not limited by Mg diffusion. However, the diffusion of Mg to the

interface to form oxides does provide vacancies that may serve as pathways for metallic A1

diffusion. As A1 diffusion in Ag is much less than that in Mg, the extent o f A1 diffusion into

the cathode, as observed by XPS, appears less for the Mg:Ag alloy than that for the Mg

alone.

Based on the fact that the chemical state o f A1 in Alq3 is A1 , the fragmentation

reaction (equation 5-1) may be modeled by an inorganic analogue such as:

2Mg + (^)A/2Oj —» 2MgO + iy^)Al (10-1)

This is a well-known oxidation-reduction reaction and is thermodynamically favored even at

room temperature. The Gibb’s free energy o f this reaction at room temperature is shown in

table 10-1. Since this metal exchange reaction also supports the behaviour o f Ag, Au, Ca [1]



and K [2,3] with Alq3 , as shown by the Gibb’s free energy values in table 10-1, metal/Alq3

interactions in general may be described by

2xM + (2/ 3)Al20 3 -a 2M xO + (%)Al (10-2)

Table 10-1 Gibb’s free energy of metal-exchange oxidation-reduction reaction at 298 K

Au Ag K Mg Ca

A G l(k J) 1660 1033 409.3 -83.6 -151.9

*Therm ochem ical data from [4]

The diffusion and reaction at the interface play a major role in the device stability over time.

It is possible that the inter-diffusion of reactive metal and reduced A1 will be significantly

slowed down only when the reacted region becomes thick enough to act as an effective

diffusion barrier. As oxides are generally also electrically insulating, the increasing thickness

o f interface oxides with time leads to the increased OLED driving voltage as a function o f

time. Should oxide growth proceed with an island type of growth pattern, dark spot formation

and eventual failure will be a likely consequence.

10.3 LiF as an interlayer

With the introduction o f LiF, the junction between the cathode and the organic is no longer

limited to a single interface. To describe the complete interfacial chemical structure, the

impact of the LiF on both the metal cathode and on the organic need to be considered.

Together, the interactions of the three components of the interface are all critical to

understanding the buried interfacial structure and the performance of the cathode in the

organic semiconductor devices.



10.3.1 L iF im pact on the cathode m eta l

10.3.1.1 Al/LiF

At thicknesses typically used in optoelectronic device cathodes, 5-10A, deposited LiF does

not completely cover the surface; instead it likely forms islands. For Al, even without

complete surface coverage, LiF is effective in decreasing the oxidation rate due to broadly

matched lattices o f the overlayer and the substrate. As LiF and Al have good lattice matching

over a broad range o f orientations, it is likely that, upon deposition, any one of the preferred

planes is aligned. The commensurate LiF islands, therefore, give the Al surface a corrugated

structure upon which the oxide grows, as in figure 10-2, with the islands acting as diffusion

barriers for Al atoms.

rnrffT

Al substrateFigure 10-2 Embedded oxide structure for oxidation of LiF coated Al surfaces.

10A LiF (-61 % coverage) is sufficient to significantly modify the oxidation kinetics,

due to an ion diffusion dominated oxidation mechanism. Ion diffusion appears to be two

orders o f magnitude faster in the oxide alone compared to the combination of LiF and oxide

on the metal surface.

When the Al is deposited on top of the LiF on an organic, such as in device

structures, the interfacial chemical structure observed is related to the protection that LiF

provides for the Al. As the thickness o f LiF increases, the protection of the cathode from

oxidation is improved. In a device, the LiF layer is mixed with an oxidized Al layer,



indicating that Al ions might be diffusing through the LiF and encountering O, either from

lateral diffusion through the organic layers or along the inorganic/organic contact. When the

metal capping layer and the interlayer have good lattice matching, the LiF layer prevents

migration of oxygen and acts as a trap for oxygen away from the metal surface. A device

with 30A LiF, for example, could last nearly 18 years on the shelf before degrading to 10%

of its initial performance when manufactured.

10.3.1.2 Mg/LiF

Deposition of LiF on Mg surfaces, which has poor nearest neighbour lattice matching, has

the opposite effect. Rather than passivating the surface, LiF on the surface changes the

products o f oxidation. Initially, there is preferential oxidation to form MgCCb on the surface,

with little apparent change in the oxide thickness. As oxidation continues, oxygen and water

likely diffuse through the incommensurate LiF lattice, and hydroxides become the dominant

oxide components. When this occurs, the oxidation rate increases rapidly, and the oxide

thicknesses for the coated and uncoated surfaces become similar. Irrespective o f the oxide

thickness, the LiF coated surfaces show preferential formation of MgC0 3 . Such carbonates

are very poorly lattice matched with Mg. The presence of LiF, therefore, modifies the

activity of the metal surface, decreasing the likelihood of Mg(OH ) 2 formation.

For Mg devices, which already show a tendency to react with the O rich groups in

organometallics, the introduction of an LiF interlayer does not protect the Mg from

destructive molecular fragmentation reactions. This suggests, along with the case of Al/LiF,

that the deposition of LiF on the organic has minimal impact on the reactivity of the organic

surface, and the interfacial reactivity can be completely described by the activity o f the metal



surface. When the two materials used for the cathode are not coherent, such as LiF with Mg,

oxidation is not prevented at the metal surface. However, the Mg/Alq3 interaction can no

longer be described by the simple analogue as before, since the introduction of LiF changes

the by-products of reaction between Mg and the organic layer. With these interfacial reaction

products, the injection o f electrons appears to be limited, which results in the complete

suppression of luminescence in devices with bi-layer cathodes. Subsequently, a device

incorporating LiF fails almost immediately, compared to devices with Mg alone.

10.3.1.3 Lattice constants as a predictive tool

Though the inorganic analogue is no longer sufficient to describe the interaction at these

interfaces, the reactivity o f the interface with LiF can be described with a simple model as

well, based on the change in the oxidation behaviour o f the metal cathode. Intuitively, one

might assume that the deposition of LiF would change the surface activity o f the organic

layer, and that this would control the oxidation behaviour of the metal deposited on the

organic surface. However, this is does not appear to be occurring, as shown by the differing

behaviour of Mg/LiF and Al/LiF cathodes. Instead, the effect o f LiF on the oxidation

behaviour o f the metal also controls the oxidation behaviour in the device as described in the

previous two sections.

The contact integrity and potential for oxidation may, therefore, be related to the

coherence o f the interfacial layers with the metal cathode. Even though these organic films

are amorphous and the inorganic films are polycrystalline, the bulk lattice constants can be

used as a rough guide in predicting the oxidation resistance and interface integrity. If the

lattices match over a broad range of orientations, the likelihood of the grains having similar,



well matched, orientations is high. Presumably, high lattice matching leads to good contacts

in devices, and high injection of charge carriers into the active layers.

Table 10-2 Lattice constant comparisons for low index planesLatticeMisfit

Best matched interface

Al/LiF 0.7%{100}//{100}, {110}//{110}, {111}//{111>

Mg/LiF 11.3% {0001 }//{l 11}

Mg/Mg(OH)21.9%

1.9%, 5%{0001}//{0001},

(ri02)/(ri02)

Mg/MgCC>3 30.8% {0001}//{0001}

L attice constants as in chapters 6 and 7

As table 10-2 shows, LiF has a good match with Al over a broad range o f lattice

planes. This may explain the improved resistance to oxidation with LiF coated surfaces, as

there are many possible orientations that will show matching, blocking surface oxidation. For

Mg, the LiF is generally poorly matched, with only a few possible matching planes. Oxygen

is, therefore, likely to penetrate to the metal surface very easily. In the long run, the affinity

o f LiF for C species may encourage the formation o f carbonate type species on the Mg

surface. Without LiF, the possibility of forming Mg(OH ) 2 is much higher. Since Mg(OFl) 2

has better matching along many orientations, it could help to explain why Mg cathodes

perform much better than Mg/LiF cathodes. In the case of Mg/LiF in devices, the

stoichiometric ratios suggest the likely formation of bulkier and more complex oxides, with

greater breakdown of the molecule than observed with Mg cathodes.



10.3.2 L iF im pact on the organic

When LiF comes into contact with electron accepting organic molecules, some o f the LiF

interacts with the conjugated carbon species, leading to formation of a charge transfer

complex. The appearance o f a high binding energy shoulder in the F Is core level can be

considered as proof of this interaction with the p bonds in the organic molecule. For C6o,

which has an abundance o f p bonds, the appearance o f the shoulder in the F Is level was

accompanied by a change in the CIs satellites for C6o similar to what has been observed

previously for chemisorption type interactions. The interaction between LiF and organic

molecules is generally complex, and cannot be described by simple inorganic analogues as

for metal/Alq3 interactions since there is no dissociation o f bonds in either the organic or in

the LiF. Additionally, there are still outstanding questions regarding the impact o f substrates

and deposition conditions on the appearance o f this interaction. There appears to be a critical

thickness for the onset o f this interaction; therefore, further work needs to be done utilizing

other techniques to clarify the growth of LiF on metal surfaces and the impact that various

substrates have on the appearance of the charge-transfer complex.

Initially, it had been speculated that the C-F interaction observed in Alq3 -LiF devices

might be a reason for the improved performance with the use of LiF interlayers, through

modification of the organic surface activity or electronic structure. However, interfacial

reactivity appears to be related more to the cathode activity than the organic, and LiF is not

always beneficial at interfaces where this interaction is visible, such as with Ag cathodes. It

is likely, therefore, that though this F-C bond is a spectroscopically observable phenomenon,

it has little real impact on the device properties.



10.3.3 L iF in terlayer properties

The formation o f the charge transfer complex is confined to the interface. As the thickness

increases, ionic LiF dominates. LiF growth does not appear to follow a layer-by-layer

mechanism, and the chemical structure and thickness o f the layer itself is the same,

irrespective of the substrate on which it is deposited. As the growth trend is very similar on

organics as on Si, it is likely that LiF growth predominantly follows an island-type growth

mode, which was also observed on metal surfaces. There appears to be minimal diffusion of

LiF into the organic layers, though the interface is roughened according to the topography of

the underlying substrate. As the thickness o f the deposited LiF layer increases, the island

coverage and island size appear to increase until the entire surface is covered. This is

occurring around 20-30A deposition. By 100A deposition, the surface is completely covered,

with the topography related to the underlying substrate.

Though island growth suggests the formation o f a complete layer on the surface after

approximately 20-30A deposition, the maximum possible useable thickness of the LiF layer

in a device is highly dependent on the nature of the underlying organic layer. There is

generally a critical thickness above which devices will no longer show adequate injection,

which is different for different organic molecules. As the LiF layer appears independent of

the underlying substrate, the conductivity differences for LiF/organic combinations observed

by X-ray photoelectron spectroscopy and in devices can only be attributed to the conductivity

o f the underlying layer. The maximum thickness for the LiF would be that at which the

potential drop across the entire system is too great to ensure adequate injection and the

device would fail. C6o with much greater conductivity than Alq3 is able to accommodate a

greater potential drop and a much thicker LiF layer is still effective in devices. The



combination of LiF/organic and the overlying metal is, therefore, analogous to a Metal-

Inorganic-Metal (MIM) type capacitor, with electron injection occurring at the “floating”

electrode (i.e. organic ETL) and emission layer interface. In such a structure, as long as the

LiF layer is thin enough to allow adequate conduction o f electrons, it can act as a charged

source of electrons that can be injected into the emission layer, like the behaviour in a flash

memory device. XPS can be used to probe the relative conductivity and dielectric properties

o f the combined dielectric/organic layer.

10.4 Metal/LiF/organic system

Based on the above observations regarding the interaction o f the metal and of the organic with

LiF, the structure for buried interfaces with Al and Mg can be summarized as in figure 10-3.

Mg

M g tQ H b iA l;-

organic

Mg

; ; ; ; ; ; ; iMgQÂI; ;.;

dgQiiS:

Figure 10-3 Schematic of interfacial structures for various metals with a LiF interlayer

organic

The interfacial structure is quite complicated and often specific to the material system

under investigation, as seen from figure 10-3. In organic electronic devices, the buried

electrode contact cannot be taken as the simple junction o f two materials. The structures



observed and the device performance can only be explained if the cathode itself is taken as a

three component system. The overlying metal layer, the LiF interlayer and the underlying

organic layer all contribute to the final interfacial structure. Changing any one component o f

the structure can have a major impact on the device performance. Often re-optimization of

the device structures is necessary in response to changes in the relative importance o f any

one factor at the interface. Therefore, the metal/LiF/organic system should really be

considered as a tri-layer cathode, with the organic layer near the interface as critical to device

performance as the other two components.

10.5 Cathode selection for organic electronics

Since the cathode structure can only be adequately explained by examining all three

components, suitability of a cathode to a particular organic depends on the relative

importance of the different interactions possible at the interfaces, as described above. For

example, in some cases, the oxidation characteristics o f the interface are more important than

the formation o f charge transfer compounds. An excellent example is the behaviour o f LiF

with Mg as a cathode. In the presence of LiF, Mg shows enhanced carbonate formation. In a

device, the introduction of LiF changes the products of reaction between Mg and the organic

active layers. Likely due to these new reaction products, the injection o f electrons is limited

and luminescence is completely suppressed with bi-layer cathodes.

The change in the surface activity with LiF-metal interactions has a greater impact

than the organic-LiF interaction in determining the ability o f the interlayer in protecting the

interface from oxidation. In contrast to the impact on Mg, LiF suppresses oxidation o f Al.

Over time, the performance of an OLED is related to the oxidation o f the cathode at the



interface with the organic layers. Therefore, increasing the thickness o f the LiF layer in a

device with an Al cathode can increase the shelf time, and maintain better device

performance over time by preventing interfacial oxidation.

As the LiF thickness increases, however, the initial device performance is often

affected. There needs to be a compromise between protecting the cathode from oxidation and

achieving optimal device performance during cathode selection. By considering the cathode

as a tri-layer structure incorporating the properties of the metal, the dielectric and the

underlying organic, more robust devices can be made using much thicker layers o f LiF with

C6o as the electron transport layer compared with what is possible for Alq3 , without greatly

sacrificing the initial device properties.

Ultimately, one o f the critical insights gained using the combined approach outlined

in this dissertation is that a universal cathode analogous to ITO as a universal anode is

unlikely. However, for good device performance, the combination of cathode and organic

layer should meet certain criteria for the interfacial structure and stability. Some o f the

criteria for interfacial contact formation and, therefore, device behaviour and stability can be

estimated prior to device fabrication from simple material property information, such as the

bulk lattice constant matching, or by assuming inorganic analogues for organic molecules.

Armed with this simple approach to assessing the suitability o f various material

combinations, it may be possible to optimize devices more quickly when a new set of organic

molecules are introduced to improve device performance, and perhaps even move away from

conventional cathode structures. As one o f the goals of the organic electronics industry is to

have all device components made from highly flexible materials, the outcomes from this

dissertation give a better description of the interfacial conditions that need to be examined



when selecting new cathodes such as those desired for eliminating inorganic cathodes in

future displays.

In this thesis, the focus has been on describing the interfacial structures in organic

semiconductor devices with archetypal and widely used organic molecules and cathode

materials using simple inorganic analogues and material property information. Currently,

device optimization relies primarily on modification o f the active organic layers. The results

o f this dissertation suggest that interfacial engineering can play a major role in future device

optimization. Organic/inorganic contacts are also critical in other organic electronic devices,

in biological sensors, and in catalysts; therefore, knowledge and control o f the interfacial

structure at the molecular level is a multi-disciplinary concern and has the potential for a

multi-application solution. By focussing on common OLED materials, new insights have

been gained and predictive methods developed to tailor interface performance.

10.6 Future work

The major outstanding questions from this work, related specifically to the interfacial

structure in OLEDs are:

• the mechanism behind the critical thickness for the emergence of the C-F charge

transfer interaction,

• the effect of faster deposition rates on LiF growth in ultrahigh vacuum conditions,

• the grain structure of LiF grown on different substrates,

• the nature of C6o growth on metal and organic substrates.

Low temperature diffusion studies o f Al and O through LiF would also be of great

benefit to confirm diffusion rates, since no data currently exist. In addition, the natural



evolution of the multi-faceted technique including peel-off would be to incorporate valence

band measurements, with either ultraviolet or synchrotron sources, to examine the work

function and density o f states at buried interfaces in devices.

Beyond those immediate questions for OLEDs specifically, building on the findings

of this thesis, some interesting avenues o f research with regards to interfacial engineering

more broadly may be related to:

• the impact o f thin dielectric coatings on the reactivity o f metal surfaces for site

specific tissue scaffolding, biosensors, and nanocatalysis,

• the need for interfacial structure matching at disordered interfaces as a

fundamental component o f interfacial engineering

• the potential extension o f X-ray photoelectron spectroscopy as a non-contact tool

for electrical measurements, and

• the potential for engineering surfaces and interfaces for the next generation

molecular devices by controlling the transition between ordered and disordered

states of matter.

One of the most intriguing outcomes o f the research to date has been to indicate that

though the organic layers are amorphous and evaporated metal layers are generally

polycrystalline, the lattice coherence between the layers appears to play an important role in

contact formation and in surface activity. This is not unusual for highly ordered systems.

Forlsch et al. [5], for example, have already shown that kinked metal surfaces can be used to

grow epitaxially constrained NaCl crystals and that these modify the surface chemical

behaviour of Cu. However, it appears that even for systems that are considered disordered,

the concept of lattice matching, as in highly ordered inorganic semiconductors, is still a good

indication of interface formation. One excellent avenue of research therefore lies in



confirming the possibility of lattice coherence as a significant factor in surface activity for

dielectrics on metal surfaces.

In order to quantify such surface effects, investigations would primarily use scanning

tunnelling (STM) and atomic force microscopies (AFM). Extension o f this research to

investigate the absorption properties o f fullerenes and phthalocyanines on the coated and

uncoated surfaces would also be o f great practical value. Annealing studies o f these surfaces

after submonolayer deposition, especially, could examine the evolution o f the organic surface

morphology, giving information about the utility o f such coated surfaces for catalytic

applications. For a metal-dielectric-organic layer, AFM could also be used in the conductive

mode for direct localized transport measurements [6]. Combined with a chemical

characterization method, such as photoelectron spectroscopy or secondary ion mass

spectroscopy, these studies would allow for a complete description of the interfacial reaction

products and electronic states that form during the absorption process.

The proposed outcome would be the prediction of effective and potentially stable

metal/di electric and organic combinations derived from an analysis o f the surface energy

constraints on growth and thin film formation. The optimal thickness o f the dielectric for

either electronic or sensing applications could then be determined from the surface structure,

rather than inferred experimentally by building prototypical devices and measuring the

current-voltage data. As the interface formation process is better understood, the significant

morphological factors for successful metal/di electric combinations may also be used to

predict novel electrode materials for the next generation of organic electronic devices.

Another potential avenue is related to the further use of XPS for non-contact

electrical measurements. It is difficult to accurately measure the conductivity o f thin organic



films and molecular devices using traditional conductance techniques, even localized ones

such as AFM. Generally, the introduction of any “external” contact to the film of interest to

probe the properties will effectively change those properties such that isolating the electrical

behaviour of the film becomes very difficult. XPS, due to its high sensitivity to the

conductivity of the films, can be used as a non-contact method for analyzing the

resistance/capacitance and other electronic properties of thin semiconductor and dielectric

films. The parameters of such a technique are not fully established but the potential for in-

si tu conduction information makes this an interesting area to explore for molecular scale

capacitors, which appear to play a large role in OLED performance [7].

There is also great potential for engineering surfaces and interfaces by controlling the

transition between ordered and disordered states. Specifically for organic electronics, there is

an intriguing phenomenon that organic transistors require a high degree o f order, whereas

organic light-emitting devices require a high degree of disorder for adequate performance.

With molecular manipulation, it is possible that one could produce a transistor/LED hybrid

device by grading the degree of disorder. Order/disorder transitions also become increasingly

important at the nano/molecular scale on surfaces, where modification of local order can

control the selective growth of self-assembled monolayers, and quantum dots. The most

likely applications of modification o f local order could include biologically interesting

sensors, which rely on specific receptor availability to trigger sensing, or site specific

nanocatalysis. Investigation of the impact of localized order on growth specifically would

require some basic high vacuum facilities, using e-beam mixing and controlled molecular

beam deposition to induce local disorder in films.



10.7 References

1 V. Choong, M. G. Mason, C. W. Tang, and Y. Gao, Appl. Phys. Lett. 72, 2689 (1998).

2 N. Johansson, T. Osada, S. Stafstrom, W. R. Salaneck, V. Parente, D. A. dos Santos, X. Crispin, and J. L. Bredas, J. Chem. Phys. 111, 2157 (1999).

3 T. Osada, P. Barta, N. Johansson, Th. Kugler, P. Broms, and W.R. Salaneck Synth. Met. 102, 1103 (1999).

4 Thermochemical Data o f Pure Substances, 3rd edition, edited by I. Barin (VCH Publishers, New York, 1989), Vol. 1.

5 S. Folsch, A. Riemann, J. Repp, G. Meyer, K.H. Rieder, Phys. Rev. B 6 6 161409 (2002); (b.) S. Folsch, A. Helms, A. Riemann, J. Repp, G. Meyer, K.H. Rieder, Surf. Sci 497 113 (2002).

6 see for example Y. Ekinci, J.P. Toennies, Surf. Sci. 563 127 (2004).

7 A. Turak, D. Grozea, Z. H. Lu, J. Elect. Spect. Rel. Phen. in prep.


Appendix A

List of empirically derived charge-binding energy correlations

A-229


A-230

Core level

C Is

for aliphatic sp3 C

Formonocyclicaromatics

for polycyclic aromatics

O ls

N Is

F I s

B i s

Ge 2p

Empirical BE vs charge

Eb (C Is) =8.0qc (a.u.) +286.2 (eV)

Eb (C Is) =4.68<?c (a.u.) +286.2 (eV) Eb (C Is) =31.06qc (a.u.) +V+0.47

Eh{C ls)=24.1qMS+29\.l (eV)

Eb(C ls)=26.9qMS+2U.2 (eV)

AEb(C ls)=25.0qMg±l.9 (eV)

Charge Investigator calculation

method

Sleigh et al. [1]

Folkesson et al. [2] Jolly and Perry [3]

ab initio AMI Mulliken population analysis

F olkesson-LarssonJolly and Perry

Grey and Hercules [4] Modified Sanderson

Sastry et al. [5] Modified Sanderson

Grey and Hercules Modified Sanderson

Eb(C 7s ) = 16.7#m s+286.8 (eV)

AEb(C ls)=l9.lqMS+0A (eV)

Eb(C 7s)=l 8 .5^ + 286 .3 (eV)

Patil et al. [6] Modified Sanderson

Grey and Hercules Modified Sanderson

Patil et al. [7] Modified Sanderson

Eb (O Is) =18.0q0 (a.u.) +538.5 (eV)

Eb (O Is) =4.23q0 (a.u.) +534.1 (eV) Eh (O Is) =30.43ô (a.u.) +V-0.27

AEb(0 Is) =n.9qMs -0.1 (eV)

Eb (N Is) =\4.\qN(a.u.) +404.9 (eV)

Eb (N Is) =7.0qN (a.u.) +401.4 (eV) Eb (N Is) =7.0qN (a.u.) +V-0.46 Eb (N'3 Is) = 2 7 .2 ^ + 4 0 3 .6 (eV)Eb (N+5 Is) =23.2qMS +405.5 (eV)

Eb (F Is) =19.6^ (a.u.) +691.7 (eV)

Eb (F Is) =4 .28^ (a.u.) +688.8 (eV) Eb (F Is) =4.28gF(a.u.) +V+1.08

AEb{F 7s)=1 1 .0<7Ms+1 .0 (eV)

ab initio AMI Mulliken Sleigh et al. , .. , .° population analysisFolkesson et al. Folkesson-Larsson

Jolly and Perry Jolly and PerryGrey and Hercules Modified Sanderson

Sleigh et al.

Folkesson et al. [8] Jolly and Perry

Grey and Hercules Grey and Hercules

Sleigh et al.

Folkesson et al. Jolly and Perry

Grey and Hercules

ab initio AMI Mulliken population analysis Folkesson-Larsson

Jolly and Perry Modified Sanderson

Modified Sanderson


Jolly and Perry

Modified Sanderson

Eb{B 7s) = 1 7 .6 ^ -2 6 (eV) Grey and Hercules Modified Sanderson

Eb (Ge 2p) =10.5flus +129.4 (eV) Grey and Hercules Modified Sanderson


A-231

S2p

Si 2p

Sn 3d5 / 2

C l 3p3/2

Br 4 p 3 /2 *

P 2 p

Ti 2p3/2

Cr 2 p 3 / 2

Rh 3d5 / 2

Pd 3d5/2

Zr 3d5 / 2

Ni 2p3/ 2

Cu 2 p 3 / 2

Eb (S2p) = 3 .3% (a.u.) +163.8(eV)

Eb (S'2 2p) =11.6<7ms+164.2 (eV)Eb (S+4 2p) =25.4qm +161.3 (eV)

Eb (S+6 2p) =20.OqMs +167.1 (eV)

Folkesson et al. Grey and Flercules

Grey and Hercules

Grey and Hercules

Folkesson-Larsson

Modified Sanderson

Modified Sanderson

Modified Sanderson

Eb (Si 2p) =1.53qsi (a.u.) +100.6(eV) Eb (Si 2 p) =11 AqMs +98.3 (eV)

Folkesson et al. Folkesson-LarssonGrey and Hercules Modified Sanderson

Eb (Sn 3d5/2) =9.0qm +491.8 (eV) Grey and Hercules Modified Sanderson

Eh (Cl 3p3/2) =6.3qa (a.u.) +201.0 (eV)

Eh (Cl 3p3/2) =4.25qc, (a.u.) +201.2 (eV) Eb (Cl 3p3/2) =6.2qMS +207.2 (eV)

, ab mitio AM I MullikenSleigh et al. , .. , .° population analysis

Folkesson et al. Folkesson-LarssonGrey and Hercules Modified Sanderson

Eb (Br 4p3/2) =5.6qMS +476.8 (eV) Grey and Hercules Modified Sanderson

Eb (P 2p) =1.67qP (a.u.) +131,6(eV)

Eb (P+3 2 p) = 1 2 .8 ^ + 1 3 6 .3 (eV)Eb (P+5 2p) =17.Oq ms +135.9 (eV)

Folkesson et al. Folkesson-LarssonGrey and Hercules Modified SandersonGrey and Hercules Modified Sanderson

Eb (Ti 2 p3/2) =5.0qn (a.u.) +454.0 (eV) Sleigh et al. ab initio AMI Mulliken population analysis

Eb (Cr 2p3/2) =4.2qCr (a.u.) +574.2 (eV) Sleigh et al.

Eb (Cr 2p3/2) =2.33qCr (a.u.) +575.3 (eV) Folkesson et al.


Eb (Rh 3ds/2) =2.5qm (a.u.) +307.3 (eV) Sleigh et al. ab initio AM I Mulliken population analysis

Eb (Pd 3dS/2) =4.1 qPli (a.u.) +335.1 (eV) Sleigh et al.

Eb (Pd 3dS/2) =4.45<7/y (a.u.) +333.9 (eV) Folkesson et al.

ab initio AM 1 Mulliken population analysis Folkesson-Larsson

Eb (Zr 3ds/2) =4.4qZr (a.u.) +178.8 (eV) Sleigh et al. ab initio AM I Mulliken population analysis

Eb (Ni 2p3/2) =6.14qNi (a.u.) +848.3 (eV) Folkesson et al. Folkesson-Larsson

Eb (Cu 2p3/2) =1.52qcu (a.u.) +932.2 (eV) Folkesson et al. Folkesson-Larsson


A-232

Pt 4 f Eb {Pd 4f) =3.\7qPl (a.u.) +71.1 (eV) Folkesson et al. Folkesson-Larsson

Mo 3d5 /2 Eb {Mo 3ds/2) =5.54qMo (a.u.) +228.2 (eV) Folkesson et al. Folkesson-Larsson

Fe 2p3/2 Eb {Fe 2p3/2) =6 AqFe (a.u.) +704.1 (eV) Folkesson et al. Folkesson-Larsson

Sn 3d5/2 Eb {Sn 3d5/2) =1.81 qFe (a.u.) +485.8 (eV) Folkesson et al. Folkesson-Larsson

Ar 3p3/2 Eb {Ar 3p3/2) =16.8qm +142.5 (eV) Grey and Hercules Modified Sanderson

Se 3d Eb {Se2 3d) =53.3qMS +60.5 (eV)Eb {Se+4 3d) =19.2qMS +58.4 (eV)Eb {S e 6 3d) =6.5 qMS +59.6 (eV)

Modified Sanderson method yields poor approximation of binding energy

Grey and Hercules

Grey and Hercules

Grey and Hercules

Modified Sanderson Modified Sanderson

Modified Sanderson

A.l References

1 C. Sleigh, A. P. Pijpers, A. Jaspers, B. Coussens, and R. J. Meier, J. Electron Spectrosc. Relat. Phenom. 77, 41 (1996).

2 B. Folkesson and R. Larsson, J. Electron Spectrosc. Relat. Phenom. 50, 267 (1990).

3 W. L. Jolly and W. B. Perry, J. Am. Chem. Soc. 95, 5542 (1973).

4 R.C. Gray and D.M. Hercules, J. Electron Spectrosc. Relat. Phenom. 12 37 (1977).

5 M. Sastry and P. Ganguly, J. Phys. Chem. A 102 697 (1998).

6 V. Patil, S. Oke, and M. Sastry, J. Electron Spectrosc. Relat. Phenom. 85 249 (1997).

7 V. Patil and M. Sastry, J. Electron Spectrosc. Relat. Phenom. 94 17 (1998).g

B. Folkesson and R. Larsson, J. Electron Spectrosc. Relat. Phenom. 50, 251 (1990).


Appendix B

Schematic of OMAC chamber

B-233


B-234


Appendix C

Data analysis in XPS

C.l Curve fitting routines

For quantitative analysis of the binding energy and shape of the core level peaks, PHI

MultiPak 6.1 A was used for least-squares analysis.

For symmetric peaks, the fitting [1] used a summation Voigt formula [2], which is a

summation of Gaussian and Lorentzian functions commonly used for core level analysis [3]:

i -% g- ln (2 )

% G *e [ FWHM +, \ 2 ( X - P P ) ¥

L FW H M J(C-l)

where A, is the binding energy value for data point i, PP is the binding energy of the peak’s center, H is the height of the peak at its center, FWHM is the full width at half maximum of the peak, %G is the percentage Gaussian component (where 0 is 0% and 1.0 is 100%)

C-235


C-236

For metals, the asymmetry on the high binding energy side of the peak must be

2 X . - P Pln(2)

FW H Maccommodated. If we designate the exponential function, H * e 1 , in the Voigt

formalism above as G, for Gaussian, the Asymmetric Voigt function can be defined as [1]

A GL(X, ) = G L(Xi) + T s ( l - — \ l * (C-2 )V H J

where TL is the tail length in half width at half maximum of the peak, and TS is the tail scale factor.

The shape of the exponential tail is defined by the TL parameter. The TS parameter is a

scaling parameter to properly size the tail to the symmetric portion of the curve.

An asymmetry parameter can therefore be defined as12

fi = T S * e TL (C' 3)

C.2 Shirley background

A Shirley background can be used to eliminate the contribution to the data from the

scattering of low energy electrons. Scattering causes an increase in the intensity o f the

emitted photoelectron on the high binding energy side, yielding a stepped appearance to the

spectrum. In Multipak, the background is defined by a right-to-left integration between two

endpoints within the original data, generating an integrated background curve [1]. For

iterated Shirley backgrounds, the Shirley fitting routine is performed five times successively,

with each iteration using the previous iteration’s background. If a sample has multiple

oxidation states for a given core level, use o f the Shirley background will always

underestimate the area of the higher binding energy components, by changing the

background value along the energy scale [4].


C-237

C.3 Determination of Al 2p binding energy for Alq3

In all peel-off samples involving Alq3 , some Al was observed at the delaminated interface.

Compositional analysis of the Ag/Alq3 interface on both the cathode and organic sides shows

stoichiometric agreement of N to Al consistent with that o f Alq3 , indicating that the Al 2p

core level observed on the organic side is that only from Al in Alq3 . Since the Alq3

component of the Al 2p peak was evident in all cases, all XPS spectra were internally

referenced to this Al 2p core level. Though the organic side o f the interface was used for

internal reference, the Au/Alq3 system showed the greatest difficulty for alignment on

analysis of the cathode side of the interface, as the strong signal from the Au 5pia core level,

expected at 74 eV, obscures the Al 2p core level signal1. To examine Al on the cathode side,

therefore, the Al 2s core level was observed instead. Using the Au 4fm core level, set at 84.0

eV [5] to account for any charging effects, and the difference in the Al 2s and Al 2p core

levels, as seen on the Ag cathode, the binding energy o f Al 2p in Alq3 was found to be 74.4

eV. All the other XPS spectra were referenced to a binding energy of 74.4 eV for the Alq3

component of the Al 2p core level. This binding energy is consistent both with reported

values o f Al in Alq3 [6 ] and with the Al 2p core level on the organic side o f the interface for

all the other cases. Since the Fermi energy of the Au 4 f core levels are well defined, charging

effects can be well accounted for by using the metallic Au spectra. Therefore, the binding

energy of Al in Alq3 can be further delineated. Assuming a spin-orbit splitting ratio of 2:1 [4],

and a peak separation of 0.45 eV [7] between 2p$n and 2pm, the Al 2pm core level was

determined, through peak deconvolution, to have a binding energy of 74.2 eV for Al in Alq3 .

1 Sputter depth profile measurements into the Au cathode confirm that the signal observed is a result of the cathode and not due to Al.


C-238

C.4 Valence band analysis for analyser resolution and calibration

Valence band XPS is very similar in nature to traditional XPS, using an intense X-ray energy

source to excite electrons from the surface region. As with core level electrons, the kinetic

energy o f the ejected electrons can be measured to determine its orbital level within the

electronic structure. In traditional XPS, the binding energy is the most important indication

of the chemical nature of the element, and all core levels have a similar shape dictated by the

probability curve for discrete electrons. In the valence band, however, the shape of the curve

is determined by the complexity o f overlapping curves close together in energy from all the

constituent elements present in the sample. The relation of the binding energy to the Fermi

level, while still important, becomes less characteristic of an element or compound than the

shape o f the curve. It is this shape of the valence band measured by XPS that has a unique

direct correlation with the density o f states (DOS) obtained by the one electron band structure

calculations [8 ]. The change in the characteristic shape of the DOS therefore can indicate

chemical or other electron exchange events occurring at the Fermi surface.

XPS for valence band measurements focuses on the energy region around the Fermi

level. As in traditional XPS, the binding energies o f the electrons are measured in reference

to the Ef. , which is by definition set at zero. If the material of interest is a metal, the Ef should

be less than the valence band maximum (VBM) and should correspond to its position on the

Fermi-Dirac (F-D) distribution function. For all real spectra, however, the Fermi edge is not

as sharp as that expected for the room temperature F-D function due to the spectrometer

resolution [9], For this thesis, the spectrometer resolution was calculated to be 0.3eV, as

determined by the width of the Fermi level, as shown in figure C-l.


C-239

Au \VB

"O 0.66

2.0 1.5 1.0 0.5 0.0 -0.5 -1.0 -1.5 -2.0Binding Energy (eV)

Figure C-l Determination of the spectrometer resolution

If the material has a band gap, however, the Ef will be above the VBM. The VBM, therefore,

can be determined as the intersection of two lines, one on the curve and one on the

background. To account for the analyzer resolution, the intersection point should be the

projected value at the midpoint of the valence band edge, as shown in figure C-2.

ZnO VB

_QCO

>>-4—*'</>c<D

3.8e'

cT3CDNCOEoz

-2.0 -4.0 - 6.06.0 4.0 2.0 0.0Binding Energy (eV)

Figure C-2 Determination of the valence band maximum (VBM) for ZnO

The distance from the VBM to the Fermi energy will correspond to half the band gap of the

material if it is undoped, and can be used to determine the type doping. However, surface

band bending can often mask this effect making it difficult to use XPS for doping

determination.


C-240

C.5 References

1 Physical Electronics, Operator's MultiPak Software Manual (Physical Electronics Inc., Eden Prairie, MN, 2000), Vol. 6 .

2 W. Voigt. Munch. Ber. 1912, 603 (1912).

3 P. M. A. Sherwood, in Practical Surface Analysis, 2nd edition, edited by D. Briggs and M. P. Seah (Wiley & Sons Ltd, New York, 1990), Vol. 1, App. 3, p.573.

4 J. E. Castle and A. M. Salvi, J. Vac. Sci. Technol. A 19, 1170 (2001).

5 J. F. Moulder, W. F. Stickle, P.E. Sobol, K.D. Bomben, Handbook o f X-ray Photoelectron Spectroscopy, edited by J. Chastain, and R.C. King, Jr. (Physical Electronics Inc., Eden Park, MN, 1995)

6 W. Song, S. K. So, J. Moulder, Y. Qiu, Y. Zhu, L. Cao, Surf. Interface Anal. 32, 70 (2001); (b.) T. P. Nguyen, J. Ip, P. Jolinat, and P. Destruel, Appl. Surf. Sci. 172, 75 (2001).

7 M. Watanabe, T. Kinoshita, A. Kakizaki, and T. Ishii, J. Phys. Soc. JPN. 6 5 ,1730 (1996).

8 T. Barr, Modern ESCA: the principles and practice o f X-ray photoelectron spectroscopy (CRC Press Inc., Boca Raton, FL, 1994).

9 P. G. Schroeder, W. N. Nelson, B. A. Parkinson, and R. Schalf, Surf. Sci. 459, 349 (2000).


Appendix D

Equations for quantitative XPS analysis

D. 1 Effective Attenuation Length

The wide use o f X-ray photoelectron spectroscopy has stimulated development o f theoretical

models to describe the transport of electrons through solids in the range o f energies typical

for X-ray excitation. The elastic and inelastic interactions of Auger electrons and

photoelectrons with atoms in the surface region are reflected in the shape of the energy

spectra [1], The best description of this transport is the effective attenuation length (EAL), a

definition of the opacity of the material for a given electron energy. Consistent and accurate

values for this parameter are necessary for any meaningful quantitative analysis from XPS

measurements. Though it is possible to measure the EAL from overlayer experiments,

analytical descriptions have been developed to aid the practical user. EALs are calculated

D-241


D-242

from analytical expressions derived from solutions o f the kinetic Boltzmann equation within

the transport approximation [2]. The EALs depend on two material-dependent parameters,

the inelastic mean free path (IMFP) and the transport mean free path (TMFP). For a complete

review o f the meaning o f the EAL and its experimental and theoretical derivation, see “The

electron attenuation length revisted” A. Jablonski and C. J. Powell, Surf. Sci. Rep. 47 33

(2002) [3]. Currently, a database is available that both contains many known EALs for

common compounds, and allows the calculation of the EAL for any material for which there

are no know values, known as the National Institute o f Standards and Technology EAL

Database [4],

D. 1.1 Inelastic Mean Free Path

The IMFP is the average distance, measured along trajectories, that particles with a given

energy will travel between inelastic collisions. A comprehensive overview of the

measurement and calculation of IMPF for a number of elements and compounds is given in

C. J. Powell and A. Jablonski, J. Phys. Chem. Ref. Data 28, 19 (1999) [5], Though a number

o f theoretical descriptions have been proposed, the most widely accepted method of

determining the IMFP in materials for which no measurements have been made is the TPP-

2M equation of Tanuma, Powell and Penn [6 ],

IM FP(D-l)

Where

0.944(D-2)

y = 0.191p~05 (D-3)


D-243

C = 1.97 -0 .9 \U (D-4)

D = 53.4-20.817 (D-5)

(D-6 )M 829.4

•3and Ep is the free-electron plasmon energy (in eV), p is the density (in g/cm ), Nv is thenumber of valence electrons per atom or molecule, M is the atomic weight, and Eg is the bnadgap energy (in eV).

Many tabulations o f these values exist, including the National Institute o f Standards

and Technology Electron IMFP Database [7], which also allows calculation of IMFP for

those materials for which no data currently exists.

D. 1.2 Transport Mean Free Path

The IMFP, however, only takes into account the inelastic scattering o f electron during

transport through the solid. The values of the practical EALs can differ from the IMFP by up

to 35% for common XPS measurement conditions due to the effects o f elastic-electron

scattering. The complete analytical description o f this electron movement requires an

additional parameter to account for this scattering. The TMFP is the average distance that an

electron must travel before its momentum in the initial direction o f motion is reduced, by

elastic scattering alone, to 1/e o f its initial value. The TMFP can be defined as [8 ]

n components , and <7tr is the transport cross section for an electron of a given energy.

Tabulations of the transport cross-section are available for most elements, and can be

determined for any element using the National Institute of Standards and Technology

where N is the atomic density, x is the atom fraction of the kth component o f a compound with


D-244

Electron Elastic-Scattering Cross-Sections Database [9], A review of the transport

approximation solutions for determination o f the cross sections is given in A. Jablonski,

Phys. Rev. B 58, 16470 (1998) [1],

D.2 Secondary Electron Yield

Following the analytical method o f Henke et al. [10], the yield o f all electrons excited by X-

ray illumination can be described by

s in 0 J0 2 k ( x1+ p/ +

^ dE‘ (D-9)

where (f) is the incident beam angle, E ,t is the kinetic energy o f all electrons, Em is the energy of the emitted electron, where Xe i s the eliminating pair mean free path, and Xp is the transport mean free path and M(E0) is defined by

M (E 0) = 4ttEofl(E11) f(E 0)PB (D-10)

where E0 is the photon energy, ju(E0) is the mass photoionization cross section at E0, p is the density, f(E 0) is the effective fraction o f absorbed photon energy lost to fluorescence and primary radiation and B is the energy loss rate for recombining electrons

For a fixed incident beam, the secondary yield can be described by

dSx _ K M (E 0) ( udEk sin ^ 2k ( X / V (X +

1+ p/ +(D-11)

~ p /

, , k )where EA is the electron affinity o f the surface, and Ek is the kinetic energy of the emitted electron.

Over the whole energy range, the yield as a function of the surface voltage, therefore,

can be given as


D-245

<?1 C =M{ E0)________ (XeApf 2

^{EA+Vs)sm<l) f %1+ V I +

(D-12)

The pair eliminating mean free path is related to the EAL by

+ (D-13)

In most cases for semiconductors, where the transport free path is very long, the two can

be taken as equivalent

D.3 References

1 A. Jablonski, Phys. Rev. B 58, 16470 (1998).

2 I. S. Tillin, A. Jablonski, J. Zemek, and S. Hucek, J. Electron Spectrosc. Relat. Phenom. 87, 127 (1997).

3 A. Jablonski and C. J. Powell, Surf. Sci. Rep. 47 33 (2002).

4 C. J. Powell and A. Jablonski, NIST Electron Effective-Attenuation-Length Database - Version 1.0, National Institute of Standards and Technology, Gaithersburg, MD, (2001).

5 C. J. Powell and A. Jablonski, J. Phys. Chem. Ref. Data 28, 19 (1999).

6 S. Tanuma, C. J. Powell, and D. R. Penn, Surf. Interface Anal. 21, 165 (1994).

7 C. J. Powell and A. Jablonski, NIST Electron Inelastic-Mean-Free-Path Database - Version 1.1, National Institute of Standards and Technology, Gaithersburg, MD, (2000).

8 A. Jablonski, Phys. Rev. B 58, 16470 (1998).

9 A. Jablonski, F. Salvat and C. J. Powell, NIST Electron Elastic-Scattering Cross-Section Database - Version 3.0, National Institute o f Standards and Technology, Gaithersburg, MD, (2002).

1 0 B. L. Henke, J. Liesegang, and S. D. Smith, Phys. Rev. B 19, 3001 (1979).


Appendix E

Structure calculations for Ceo-LiF interaction

E.l Geometry optimized structures and theoretical prediction of the F I s core level shift

Density functional calculations are useful for determining the geometry optimized interaction

between LiF and C6 o and potentially test the validity of the semi-ionic bonding proposal.

Such calculations were performed in collaboration with Dr. Dharma-Wardana at the National

Research Council. Many possible arrangements arise for the LiF-C6 o interaction when

modeled since individual interacting molecules as the Li or the F unit may be placed close to

the hexagonal faces, pentagonal faces, or one of the C-C bonds in the C 6 o molecule. The

various arrangements for the interaction between LiF and Cgo were geometry optimized by

total energy minimization to ensure that they represent realistic structures. The Mullikan

E-246


E-247

charges and binding energies were determined for an ionic LiF molecule, for an isolated LiF

molecule, and for fullerenes interacting with a single LiF molecule in three configurations:

(1) Li close to a hexagonal face with the Li-F bond normal to the face (2) F close to the

hexagonal face and (3) Li-F bond slanted to have the F atom above a pentagonal face. A few

other cases with a fullerene molecule in contact with two LiF molecules were also

investigated. In each case, the atomic positions were optimized by total-energy minimization.

These electronic-structure details are obtained from density functional calculations using the

Gaussian-98 code [ 1 ]. The exchange-correlation effects were treated using the BP8 6

functionals of Becke et al. [2] where Gaussian basis functions o f the 6-31G* type were used.

(For acronyms, basis sets, etc. see [2] and [3]). The geometry optimized bond lengths,

Mullikan charges and binding energies are summarized for all the cases in Table E - l.

Table E-l Theoretical bond lengths, binding energies assuming Koopman’s approximation and Mullikan charges binding energy calculation for model structures of LiF-C60 interaction______________________________________

Structure Q F <lu qcJ (-60 r FLi (A) Eb F Is

LiF -0.5 +0.5 — 1.586 656.304

Ceo-FLi1 -0.491 +0.523 -0.0365 1.573 657.359

C6 0 -LiF2 -0.5 +0.365 +0.135 1.574 656.348

C6 0 -LiF3 -0.499 +0.348 +0.153 1.577 656.448

LiF-C6 0 -LiF4 -0.493 +0.535 +0.042 1.573 657.645

-0.504 +0.347 +0.157 1.574 656.100

LiF-C6 0 -FLi5 -0.490 +0.524 +0.034 1.573 657.151

-0.489 +0.521 +0.032 1.573 657.126

FLi-Ceo-LiF6 -0.480 +0.374 +0.106 1.572 656.611

-0.489 +0.355 +0.134 1.573 656.6101F near a hexagonal face, the LiF bond is normal to the face.2Li near a hexagonal face, the LiF bond is nonnal to the face.3Li is on a bond between a hexagonal and a pentagonal face. The LiF bond is slanted so that the F atom is

above the pentagonal face.4 F near a hexagonal face, Li near the opposite hexagonal face. LiF bonds nonnal to the hexagonal faces5 same as above but an F is adjeacent to a hexagonal and its opposite hexagonal face as well6 same as above but with Li adjacent to both hexagonal faces


E-248

Contrary to expectations, this geometry optimized model suggests that the charge

distribution and bond length o f the LiF molecule in the LiF-C6 o interaction resemble that o f a

covalent LiF molecule. The charge redistribution and shortening o f the bond would also be

expected to affect the XPS spectrum by producing a high energy shoulder, as was observed.

However, the model focussed on individual molecule-molecule interactions, neglecting

relaxation effects. Therefore, the Hartree-Fock binding energy values using Koopman’s

theorem, given in table 8.2, underestimate the measured binding energy o f solid state LiF by

~30eV. As there were also no features visible in the XPS spectrum at those binding energies,

it is likely that this covalent model description was too simplistic to adequately describe the

interaction between LiF and C6 o-

E.2 References

1 Gaussian 98, Revision A.9. M.J. Frisch, G. W. Trucks, H. B. Schlegel, G. E. Scuseria, M. A. Robb, J. R. Cheeseman, V. G. Zakrzewski, J. A. Montgomery, R. E. Stratmann, J. C. Burant, S. Dapprich, J. M. Millam, A. D. Daniels, K. N. Kudin, M. C. Strain, O. Farkas, J. Tomasi, V. Barone, M. Cossi, R. Cammi, B. Mennucci, C. Pomelli, C. Adamo, S. Clifford, J. Ochterski, G. A. Petersson, P. Y. Ayala, Q. Cui, K. Morokuma, D. K. Malick, A. D. Rabuck, K. Raghavachari, J. B. Foresman, J. Cioslowski, J. V. Ortiz, B. B. Stefanov, G. Liu, A. Liashenko, P. Piskorz, I. Komaromi, R. Gomperts, R. L. Martin, D. J. Fox, T.Keith, M. A. Al-Laham, C. Y. Peng, A. Nanayakkara, C. Gonzalez, M. Challacombe, P. M. W. Gill, B. G. Johnson, W. Chen, M. W. Wong, J. L. Andres, M. Head-Gordon, E. S. Replogle and J. A. Pople, Gaussian Inc., Pittsburgh, PA (1998)

2 A.D. Becke, J. Chem. Phys. 98, 5648 (1993).

3 C. Lee, W. Yang and R.G. Parr, Phys. Rev. B 37, 785 (1988).


Appendix F

Summary of observed F Is core level in all experiments performed and observations regarding the experimental results.

F-249


F-250

Experiment F Is Observations

Depositions:Alq3 on LiF - AA deposition series/other

LiF on Alq3 - LA/LDiff deposition series

C6 o on thick LiF - CD deposition series

Cô deposited on 5A LiF on Pt - Cl series

C6 o deposited on 5A LiF on Ag - CG series

C6 o deposited on 5A LiF on ITO - CF/CH series

C6 o/LiF multilayers - CLA series

LiF on C6 o - LC/LDiff deposition series

Single peak

Single peak

Double peak

Single peak

Double peak

Single peak

C6 o deposited on 5A LiF on Au - CE series Double peak

Double peak

Double peak

No emergence of shoulder for thickness greater than 10A

Appearance of C-LiF only after 6ML deposition

No emergence of the shoulder

Very small shoulder in F Is

No emergence of the shoulder

Appearance of shoulder only after 2ML deposition

Inconsistent growth of the C-LiF shoulder in F Is, not correlated to C Is satellites

Shoulder only apparent for thickness less than 10A LiF

LiF on TPD - LT deposition series

M deposition series on LiF/Alq3

15 A Al 15 A A g

30, 60A Al

30, 60A Ag

Single peak

Single peak Single peak

Double peak

Double peak

LiF on metal (Ag, Au, Pt, Al, Mg, Cr, ITO) Single peak

Only after a critical thickness of metal deposition does the double peak appear - sufficient to supply 1 atom per LiF molecule

If chamber walls of MAC system are hot, shoulder appears intermittently on metal surfaces


F-251

Device peel-off:T 1 device series

LiF/Al Single peak

LiF/Mg:Ag Single peak

Ag/LiF/Alq3 device series (R series) Double peak

Mg/LiF/ Alq3 device series (D series) Single peak

Cr/LiF/Alq3 device series (Cr series) Double peak

Al/LiF/C6 o device series (CJ series) Double peak

Al/3,5,15,200 A LiF/Alq3 device series (Lu series) Double peak

Ag/5A A1/5A LiF/Alq3 Double peak

Breakdown reaction seen in N Is, C IsLess reaction observed than with Al, N Is peak not as pronounced, C Is shows same effect

Penetration of Ag into the organic layer indicates incomplete coverage of Alq3 by LiF

Molecular breakdown reaction of Alq3 with Al 2p shoulder

Very weak F Is

Suppression of Alq3 breakdown with thick LiF layers


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