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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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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 C^o..............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
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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) C^o 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
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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
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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
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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
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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
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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
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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
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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
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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
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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].
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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.
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Chapter 1 Introduction 3
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,
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Chapter 1 Introduction 4
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
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Chapter 1 Introduction 5
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.
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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
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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
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Chapter 2 OLED Fundamentals 8
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
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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].
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Chapter 2 OLED Fundamentals 10
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.
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Chapter 2 OLED Fundamentals 11
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.
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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
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Chapter 2 OLED Fundamentals 13
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.
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Chapter 2 OLED Fundamentals 14
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
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Chapter 2 OLED Fundamentals 15
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
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Chapter 2 OLED Fundamentals 16
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.
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Chapter 2 OLED Fundamentals 17
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
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Chapter 2 OLED Fundamentals 18
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] ^Estimated 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].
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Chapter 2 OLED Fundamentals 19
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
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Chapter 2 OLED Fundamentals 20
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.
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Chapter 2 OLED Fundamentals 21
— 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 2 OLED Fundamentals 22
2.4 References
1 M. D’lorio. Can. J. Phys. 78, 231 (2000).
2 For a review of the history and current status of organic electroluminescence see for example U. Mitschke, and P. Bauerle, J. Mater. Chem. 10, 1471 (2000), or the reviews cited in ref [11] in Chapter 1.
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4 M. A. Baldo and S. R. Forrest, Phys. Rev. B 64, art. no. 085201 (2001).
5 See for example T.-P. Nguyen, P. Molinie, and P. Destruel, in Handbook o f Advanced Electronic Materials and Devices, edited by H. S. Nalwa (Academic Press, 2001), Vol. 10: Light-Emitting Diodes, Lithium Batteries and Polymer Devices, Chap. 1, p.l.
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Chapter 2 OLED Fundamentals 23
25B. Miller, J. M. Rosamilia, G. Dabbagh, R. Tycko, R. C. Haddon, A. J.Muller, W. Wilson, D.W. Murphy, and A. F. Hebard, J. Am. Chem. Soc. 113, 6291 (1991).
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Chapter 2 OLED Fundamentals 24
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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 -
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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.
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Chapter 3 X-ray photoelectron spectroscopy 27
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]
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Chapter 3 X-ray photoelectron spectroscopy 28
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
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Chapter 3 X-ray photoelectron spectroscopy 29
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)
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Chapter 3 X-ray photoelectron spectroscopy 30
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
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Chapter 3 X-ray photoelectron spectroscopy 31
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.
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Chapter 3 X-ray photoelectron spectroscopy 32
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,
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Chapter 3 X-ray photoelectron spectroscopy 33
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
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Chapter 3 X-ray photoelectron spectroscopy 34
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
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Chapter 3 X-ray photoelectron spectroscopy 35
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]
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Chapter 3 X-ray photoelectron spectroscopy 36
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:
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Chapter 3 X-ray photoelectron spectroscopy 37
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
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Chapter 3 X-ray photoelectron spectroscopy 38
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].
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Chapter 3 X-ray photoelectron spectroscopy 39
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
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Chapter 3 X-ray photoelectron spectroscopy 40
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
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Chapter 3 X-ray photoelectron spectroscopy 41
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.
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Chapter 3 X-ray photoelectron spectroscopy 42
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
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Chapter 3 X-ray photoelectron spectroscopy 43
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
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Chapter 3 X-ray photoelectron spectroscopy 44
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.
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Chapter 3 X-ray photoelectron spectroscopy 45
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.
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Chapter 3 X-ray photoelectron spectroscopy 46
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.
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Chapter 3 X-ray photoelectron spectroscopy 47
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
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Chapter 3 X-ray photoelectron spectroscopy 48
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].
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Chapter 3 X-ray photoelectron spectroscopy 49
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.
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Chapter 3 X-ray photoelectron spectroscopy 50
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
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Chapter 3 X-ray photoelectron spectroscopy 51
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>
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Chapter 3 X-ray photoelectron spectroscopy 52
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.
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Chapter 3 X-ray photoelectron spectroscopy 53
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
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Chapter 3 X-ray photoelectron spectroscopy 54
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.
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Chapter 3 X-ray photoelectron spectroscopy 55
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
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Chapter 3 X-ray photoelectron spectroscopy 56
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.
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Chapter 3 X-ray photoelectron spectroscopy 57
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
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Chapter 3 X-ray photoelectron spectroscopy 58
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.
3.7 References
1W. R. Salaneck, S. Stafstrom, and J.-L. Bredas, Conjugated polymer surfaces and interfaces, (Cambridge University Press, Cambridge, 1996).
S. Hufner, Photoelectron Spectroscopy, (Springer, Berlin, 1995).-3
K. Siegbahn, C. Nordling, A. Fahlman, R. Nordberg, K. Hamrin, J. Hedman, G. Johansson, T. Bergmark, S. E. Karlsson, I. Lindgren, and B. Lindberg, Nova Acta Regiae Soc. Sci., Ups., 4, 20 (1967).
4 A. Scholl, Y. Zou, M. Jung, Th. Scmidt, R. Fink, and E. Umbach, J. Chem. Phys. 121, 10260 (2004).
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Chapter 3 X-ray photoelectron spectroscopy 59
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Chapter 3 X-ray photoelectron spectroscopy 60
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Chapter 3 X-ray photoelectron spectroscopy 61
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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 -
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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
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Chapter 4 Experimental 64
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
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Chapter 4 Experimental 65
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.
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Chapter 4 Experimental 66
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
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Chapter 4 Experimental 67
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.
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Chapter 4 Experimental 68
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
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Chapter 4 Experimental 69
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.
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Chapter 4 Experimental 70
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
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Chapter 4 Experimental 71
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
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Chapter 4 Experimental 72
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
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Chapter 4 Experimental 73
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
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Chapter 4 Experimental 74
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
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Chapter 4 Experimental 75
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
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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.
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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
Chapter 4 Experimental 78
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,
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Chapter 4 Experimental 79
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).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Chapter 4 Experimental 80
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).
14 A. Ranjagopal and K. Khan, J. Appl. Phys. 84, 355 (1998).
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).
28J. McElvain, H. Atoniadias, M.R. Hueschen, J.N. Miller, D. M. Roitman, J.R. Sheats, and R.L. Moon, J. Appl. Phys. 80, 6002 (1996).
29 A. Murase, M. Ishii, S. Tokito, and Y. Taga, Anal. Chem. 73, 2245 (2001).
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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 -
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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
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Chapter 5 Metal/Alq3 interface structures 83
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.
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Chapter 5 Metal/Alq3 interface structures 84
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.
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Chapter 5 Metal/Alq3 interface structures 85
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
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Chapter 5 Metal/Alq3 interface structures 86
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)
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Chapter 5 Metal/Alq3 interface structures 87
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.
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Chapter 5 Metal/Alq3 interface structures 88
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.
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Chapter 5 Metal/Alq3 interface structures 89
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
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Chapter 5 Metal/Alq3 interface structures 90
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.
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Chapter 5 Metal/Alq3 interface structures 91
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.
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Chapter 5 Metal/Alq3 interface structures 92
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.
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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
1 C. W. Tang and S.A. VanSlyke, Appl. Phys. Lett. 51, 913 (1987).
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.
5 A. Ranjagopal and K. Khan, J. Appl. Phys. 84, 355 (1998).
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).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Chapter 5 Metal/Alq3 interface structures 94
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).
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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 -
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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.
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 97
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
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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.
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 99
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)
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 100
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.
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 101
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
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 102
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
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 103
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
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 104
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
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 105
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.
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 106
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
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1/th
ickn
ess
(1/A
)
Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 107
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.
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 108
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
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 109
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,
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 110
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
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 111
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
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 112
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
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 113
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.
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 114
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
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 115
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
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 116
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)
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 117
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
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Inte
nsity
(a
rb.
units
) In
tens
ity
(arb
. un
its)
Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 118
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.
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 119
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.
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Chapter 6 LiF/metal bilayer I: Case of Al/LiF 120
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[.
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 121
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
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 122
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
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Chapter 6 LiF/metal bilayer I: Case o f Al/LiF 123
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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).
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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 -
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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.
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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
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Chapter 7 LiF/metal bilayer II: Case of Mg/LiF 127
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)
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Chapter 7 LiF/metal bilayer II: Case of Mg/LiF 128
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
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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)
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 130
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
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 131
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
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 132
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
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 133
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
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 134
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
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Chapter 7 LiF/metal bilayer II: Case of Mg/LiF 135
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.
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 136
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
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 137
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
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 138
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.
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 139
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.
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 140
— 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
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Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 141
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).
5 N. S. McIntyre and C. Chen, Corr. Sci. 40, 1697 (1998).
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).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Chapter 7 LiF/metal bilayer II: Case o f Mg/LiF 142
9 B. R. Strohmeier, Surf. Interface Anal. 15, 51 (1990).
10 J. E. Gray and B. Luan, J. Alloys Compounds 336, 88 (2002).
11 N. S. McIntyre and C. Chen, Corr. Sci. 40, 1697 (1998).
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).
23 D. Grozea, A. Turak, X.D. Feng, Z.H. Lu, D. Johnson, R. Wood. Appl. Phys. Lett. 81, 3173 (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).
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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 -
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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
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Chapter 8 Organic/LiF interaction 145
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.
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Chapter 8 Organic/LiF interaction 146
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
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Chapter 8 Organic/LiF interaction 147
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)).
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Chapter 8 Organic/LiF interaction 148
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)
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Chapter 8 Organic/LiF interaction 149
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.
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Chapter 8 Organic/LiF interaction 150
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
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Chapter 8 Organic/LiF interaction 151
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.
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Chapter 8 Organic/LiF interaction 152
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.
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Chapter 8 Organic/LiF interaction 153
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
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Chapter 8 Organic/LiF interaction 154
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 C^o 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.
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Chapter 8 Organic/LiF interaction 155
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)
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Chapter 8 Organic/LiF interaction 156
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-
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Chapter 8 Organic/LiF interaction 157
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
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Chapter 8 Organic/LiF interaction 158
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
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Chapter 8 Organic/LiF interaction 159
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
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Chapter 8 Organic/LiF interaction 160
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.
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Chapter 8 Organic/LiF interaction 161
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.
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Chapter 8 Organic/LiF interaction 162
^ _□ —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
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Chapter 8 Organic/LiF interaction 163
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.
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Chapter 8 Organic/LiF interaction 164
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
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Chapter 8 Organic/LiF interaction 165
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.
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Chapter 8 Organic/LiF interaction 166
[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],
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Chapter 8 Organic/LiF interaction 167
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 ^A 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]
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Chapter 8 Organic/LiF interaction 168
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
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Chapter 8 Organic/LiF interaction 169
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. C^o 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
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Chapter 8 Organic/LiF interaction 170
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-
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Nor
mal
ized
In
tens
ity
(arb
. un
its)
Chapter 8 Organic/LiF interaction 171
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.
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Chapter 8 Organic/LiF interaction 172
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 C^o, 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.
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Chapter 8 Organic/LiF interaction 173
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.
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Chapter 8 Organic/LiF interaction 174
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
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Chapter 8 Organic/LiF interaction 175
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
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Chapter 8 Organic/LiF interaction 176
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,
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Chapter 8 Organic/LiF interaction 177
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.
8.3 References
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Chapter 8 Organic/LiF interaction 178
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31 N. Troullier and J. L. Martins, Phys. Rev. B 46, 1754 (1992).
32 J.A. Leiro, M.H. Heinonen, T. Laiho, and I.G. Batirev. J. Elec. Spec, and Rel. Phen, 128 205 (2003).
33 Gaussian 9 8 , 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)
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Chapter 8 Organic/LiF interaction 179
34 K. Siegbahn, C. Nordling, A. Fahlman, R. Nordberg, K. Hamrin, J. Hedman, G.Johansson, T. Bergmark, S. E. Karlsson, I. Lindgren, and B. Lindberg, Nova Acta Regiae Soc. Sci., Ups., 4,20(1967).
35 C. Sleigh, A. P. Pijpers, A. Jaspers, B. Coussens, and R. J. Meier, J. Electron Spectrosc. Relat. Phenom. 77, 41 (1996).
36 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.
37 NIST X-ray Photoelectron Spectroscopy Database - Version 3.4 (Web version) , National Institute of Standards and Technology, Gaithersburg, MD, (2003).
38 R. N. Euwema, G. G. Wepfer, G. T. Surratt, and D. L. Wilhite, Phys. Rev. B 9, 5249 (1974).
39 C. P. Baskin, J. Am. Chem. Soc. 95, 5868 (1973).
40 G. Morretti, J. Electron Spectrosc. Relat. Phenom. 95, 95 (1998).
41 G. M. Rothberg, M. Eisenstadt, and P. Kusch, J. Chem. Phys. 30, 517 (1959); (b.) M. Eisenstadt, J. Chem. Phys. 29, 797 (1958); (c.) R. F. Porter, and R. C. Schoonmaker, J. Chem. Phys. 29, 1070 (1958).
42 M. F. Butman, A. A. Smirnov, L. S. Kudin, and Z. A. Munir, J. Mater. Synth. Process. 8, 93 (2000).
43 T. M. Schaub, D. E. Biirgler, and H.-J. Giintherodt, Europhys. Lett. 36, 601 (1996).
44 M. Breitholtz, J. Algdal, T. Kihlgren, S-A. Lingren, and L. Wallden, Phys. Rev. B 30, 4761 (1984).
45 C. Enkvist, St. Lunell, B. Sjogren, S. Svensson, P. A. Briihwiler, A. Nilsson, A.J.Maxwell, and N. Martensson. Phy. Rev. B 48 14629 (1993).
46 T.R. Ohno, Y. Chen, S.E. Harvey, G.H. Kroll, J.H. Weaver, R.E. Haufler, and R.E. Smalley, Phys. Rev. B 44, 13747 (1991).
47 E. Umbach, Surf. Sci. 117, 482 (1982).
48 H. Weaver, J.L. Martins, T. Komeda, Y. Chen, T.R. Ohno, G.H. Kroll, N. Troullier, R. E. Haufler, and R.E. Smalley, Phys. Rev. Lett. 66, 1741 (1991).
49 R. C. Haddon, L. E. Brus, and K. Raghavachari, Chem. Phys. Lett. 125, 459 (1986).
50 O. Gunnarsson and K. Schunhammer. Phys. Rev. Lett. 41, 1608 (1978); (b.) K. Schunhammer and O. Gunnarsson. Solid State Commun. 23, 691 (1977)
51 See for example Q. T. Le, L. Yan, Y. Gao, M. G. Mason, D. J. Giesen, and C. W. Tang, J. Appl. Phys. 87, 375 (2000); (b.) 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); (c.) L. S. Hung, C. W. Tang, and M. G. Mason, Appl. Phys. Lett. 70, 152 (1997); (d.) T. Mori, H. Fujikawa, S. Tokito, and Y. Taga, Appl. Phys. Lett. 73, 2763 (1998).
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Chapter 8 Organic/LiF interaction 180
52 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); (b.) W.J.H. van Gennip, J.K.J. van Duren, P.C. Thiine, R.A.J. Janssen, J.W. Niemantsverdriet, J. Chem. Phys. 117, 5031 (2002).
53 V. N. Mitkin, J. Struct. Chem. 44, 82 (2003).
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55 A. I. Druzhinina, N. A. Galeva, R. M. Varushchenko, O. V. Boltalina, and L. N. Sidorov,J. Chem. Thermodynamics 31, 1469 (1999).
56 O. V. Boltalina, A. D. Darwish, J. M. Street, R. Taylor, and X-W. Wei, J. Chem. Soc., Perkin Trans. 2 2002, 251 (2002).
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68 A. Kahn, J. Hwang, A. Wan, and W. Zhao, MRS Fall Meeting 2005 Symposium I Interfaces in Organic and Molecular Electronics 16.3; (b.) A. S-C. Wan, J. Hwang, F.Amy, and A. Kahn, MRS Fall Meeting 2005 Symposium I Interfaces in Organic and Molecular Electronics 110.8
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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 -
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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
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Chapter 9 LiF interlayer properties on surfaces 183
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
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Chapter 9 LiF interlayer properties on surfaces 184
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
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Chapter 9 LiF interlayer properties on surfaces 185
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^
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Chapter 9 LiF interlayer properties on surfaces 186
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
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Chapter 9 LiF interlayer properties on surfaces 187
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
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Chapter 9 LiF interlayer properties on surfaces 188
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
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Chapter 9 LiF interlayer properties on surfaces 189
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.
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Chapter 9 LiF interlayer properties on surfaces 190
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 C^o 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.
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Chapter 9 LiF interlayer properties on surfaces 191
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
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Chapter 9 LiF interlayer properties on surfaces 192
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
Nominal thickness (A)
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
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Chapter 9 LiF interlayer properties on surfaces 193
A Ey _ chKgmg ^e
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
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Chapter 9 LiF interlayer properties on surfaces 194
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
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Chapter 9 LiF interlayer properties on surfaces 195
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
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Chapter 9 LiF interlayer properties on surfaces 196
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.
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Chapter 9 LiF interlayer properties on surfaces 197
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
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Chapter 9 LiF interlayer properties on surfaces 198
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
Nominal thickness (A)
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
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Chapter 9 LiF interlayer properties on surfaces 199
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.
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Chapter 9 LiF interlayer properties on surfaces 200
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)
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Chapter 9 LiF interlayer properties on surfaces 201
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
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Chapter 9 LiF interlayer properties on surfaces 202
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.
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Chapter 9 LiF interlayer properties on surfaces 203
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
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Chapter 9 LiF interlayer properties on surfaces 204
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.
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Chapter 9 LiF interlayer properties on surfaces 205
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.
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Chapter 9 LiF interlayer properties on surfaces 206
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.
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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]
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Chapter 9 LiF interlayer properties on surfaces 208
9.5 References
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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).
4I. G. Hill, D. Milliron, J. Schwartz, and A. Kahn, Appl. Surf. Sci. 166, 354 (2000).
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).
15 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).
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.
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Chapter 9 LiF interlayer properties on surfaces 209
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).
28 J. Cazaux, J. Elect. Spect. Rel. Phenom. 113 15 (2000).
29 J. Cazaux, J. Elect. Spect. Rel. Phenom. 105 155 (1999).
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).
34 D. Grozea, A. Turak, X.D. Feng, Z.H. Lu, D. Johnson, R. Wood. Appl. Phys. Lett. 81, 3173 (2002).
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.
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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 -
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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.
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Chapter 10 Interfacial structure models and conclusions 212
• 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
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Chapter 10 Interfacial structure models and conclusions 213
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]
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Chapter 10 Interfacial structure models and conclusions 214
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.
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Chapter 10 Interfacial structure models and conclusions 215
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,
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Chapter 10 Interfacial structure models and conclusions 216
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
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Chapter 10 Interfacial structure models and conclusions 217
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,
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Chapter 10 Interfacial structure models and conclusions 218
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.
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Chapter 10 Interfacial structure models and conclusions 219
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.
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Chapter 10 Interfacial structure models and conclusions 220
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
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Chapter 10 Interfacial structure models and conclusions 221
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
; ; ; ; ; ; ; iMgQ^AI; ;.;
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
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Chapter 10 Interfacial structure models and conclusions 222
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
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Chapter 10 Interfacial structure models and conclusions 223
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
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Chapter 10 Interfacial structure models and conclusions 224
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
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Chapter 10 Interfacial structure models and conclusions 225
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
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Chapter 10 Interfacial structure models and conclusions 226
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
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Chapter 10 Interfacial structure models and conclusions 227
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.
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Chapter 10 Interfacial structure models and conclusions 228
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.
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Appendix A
List of empirically derived charge-binding energy correlations
A-229
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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) =30.43^o (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
ab initio AMI Mulliken population analysis Folkesson-Larsson
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
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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.
ab initio AMI Mulliken population analysis Folkesson-Larsson
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
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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).
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Appendix B
Schematic of OMAC chamber
B-233
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B-234
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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
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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].
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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.
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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.
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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.
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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).
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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
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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)
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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
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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
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<?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).
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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
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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
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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).
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Appendix F
Summary of observed F Is core level in all experiments performed and observations regarding the experimental results.
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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
C^o 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
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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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