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Silicon carbide biocompatibility, surface control, and electronic cellular interaction for biosensing applications


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Silicon carbide biocompatibility, surface control, and electronic cellular interaction for biosensing applications
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Coletti, Camilla
University of South Florida
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Tampa, Fla.
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Biological cell
Hybrid system
Surface topography
Surface potential
Surface chemistry
Dissertations, Academic -- Electrical Engineering -- Doctoral -- USF   ( lcsh )
bibliography   ( marcgt )
theses   ( marcgt )
non-fiction   ( marcgt )


ABSTRACT: Cell-semiconductor hybrid systems are a potential centerpiece in the scenery of biotechnological applications. The selection and study of promising crystalline semiconductor materials for bio-sensing applications is at the basis of the development of such hybrid systems. In this work we introduce crystalline SiC as an extremely appealing material for bio-applications. For the first time we report biocompatibility studies of different SiC polytypes whose results document the biocompatibility of this material and its capability of directly interfacing cells without the need of surface functionalization. Since the successful implementation of biosensors requires a good understanding and versatile control of the semiconductor surface properties, the chemistry, crystallography and electronic status of different SiC surfaces are extensively studied while their surface morphologies are thoroughly controlled via hydrogen etching. Also, investigations of the effect of cell surface charge on the electronic status of SiC surfaces are attempted adopting a contactless surface potential monitoring technique. The results obtained from these contactless measurements lead to the development of theoretical models well-suited for the description of cell-semiconductor hybrid systems electronic interactions.
Dissertation (Ph.D.)--University of South Florida, 2007.
Includes bibliographical references.
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by Camilla Coletti.
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Includes vita.

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Silicon carbide biocompatibility, surface control, and electronic cellular interaction for biosensing applications
h [electronic resource] /
by Camilla Coletti.
[Tampa, Fla.] :
b University of South Florida,
3 520
ABSTRACT: Cell-semiconductor hybrid systems are a potential centerpiece in the scenery of biotechnological applications. The selection and study of promising crystalline semiconductor materials for bio-sensing applications is at the basis of the development of such hybrid systems. In this work we introduce crystalline SiC as an extremely appealing material for bio-applications. For the first time we report biocompatibility studies of different SiC polytypes whose results document the biocompatibility of this material and its capability of directly interfacing cells without the need of surface functionalization. Since the successful implementation of biosensors requires a good understanding and versatile control of the semiconductor surface properties, the chemistry, crystallography and electronic status of different SiC surfaces are extensively studied while their surface morphologies are thoroughly controlled via hydrogen etching. Also, investigations of the effect of cell surface charge on the electronic status of SiC surfaces are attempted adopting a contactless surface potential monitoring technique. The results obtained from these contactless measurements lead to the development of theoretical models well-suited for the description of cell-semiconductor hybrid systems electronic interactions.
Dissertation (Ph.D.)--University of South Florida, 2007.
Includes bibliographical references.
Text (Electronic dissertation) in PDF format.
System requirements: World Wide Web browser and PDF reader.
Mode of access: World Wide Web.
Title from PDF of title page.
Document formatted into pages; contains 180 pages.
Includes vita.
Advisor: Stephen E. Saddow, Ph.D.
Biological cell.
Hybrid system.
Surface topography.
Surface potential.
Surface chemistry.
Dissertations, Academic
x Electrical Engineering
t USF Electronic Theses and Dissertations.
4 856


Silicon Carbide Biocompatibility, Surface Control and Electronic Cellular Interaction for Biosensing Applications by Camilla Coletti A dissertation submitted in partial fulfillment of the requirements for the degree of Doctor of Philosophy in Electrical Engineering Department of Electrical Engineering College of Engineering University of South Florida Major Professor: Stephen E. Saddow, Ph.D. Andrew M. Hoff, Ph.D. Mark J. Jaroszeski, Ph.D. Ulrich Starke, Ph. D. Jing Wang, Ph. D. Date of Approval: October 9, 2007 Keywords: biological cell, semiconductor, hybrid system, surface topography, surface potential, surface chemistry Copyright 2007, Camilla Coletti


Dedication To Gene, for making home a better place, for having the patience to be with me 24/7, and for sharing with me the good and th e bad of most of these three years of ‘Florida-life’ (including the long sleepless nights which re sulted in this manuscript). To Mom Anna Enrica, Dad Fabiano and Sis Giulia for all the love, the continuous support and the daily endless phon e calls which made me feel at home in Italy with them every day of these three l ong years. Despite the ocean.


Acknowledgments In the first place I wish to thank Dr. S. E. Saddow, who offered me the chance to start a doctorate prog ram in his group and provided mo ral and financial support through all these years. My gratitude also goes to Dr. A.M. Hoff, Dr. M.J. Jaroszeski, Dr. U. Starke, and Dr. J. Wang, for agreeing to serve on my Ph.D. committee and for providing valuable suggestions for the improvement of the manuscript. For their assistance in LEED and AES measurements (Chapter 2), my gratitude goes to M. Hetzel and Dr. C. Virojanadara of Dr. Starke’s group at the MPI of Stuttgart (DE). C. Coggins of PSIA (Santa Clara, CA) is ackowledeged for her contribution in AFM imaging (§ 3.1.4). I wish to thank Dr. Iannotta of the IFN-CNR of Tr ento (IT) and his group, in particular A. Pallaoro and T. Toccoli for their contribution in the wettability studies and Drs. M. Nardi and R. Verrucchi for their contribution in XP S studies (§ 3.2.3 and § 3.3.3). I am thankful to Dr. B. Grayson and J. Jimenez of USF fo r their valuable assist ance in XPS and ATRFTIR measurements, respectivel y (§ 4.4.2). I owe the survival of my cell cultures in my periods of absence from the lab to R. C onnelly, N. Ramachandran and D. Bazley to whom I am most grateful. For technical s upport with CPD instru mentation I wish to thank Dr. E. Oborina and SDI (Semiconducto r Diagnostic Inc., Tampa, FL). For providing a very enjoyable working atmosphe re and helping me in many issues whose list would be too long to be contained in one page I deeply th ank my colleagues and friends: Chris Frewin, Ian Haselbarth, Chris Locke, Norelli Schettini, and Dr. Jeremy Walker.


i Table of Contents List of Tables iv List of Figures vi Abstract xii Chapter 1. Introduction 1 1.1. Research objective and motivation 1 1.2. Silicon carbide: a promising material for bio-sensing applications 5 1.2.1. SiC general properties 5 1.2.2. SiC as a biomaterial: background information 7 1.3. Surface characterization tools 9 1.4. Contact potential difference t echnique for cell-semiconductor electronic interaction studies 13 1.4.1. Semiconductor energy band diagrams 14 1.4.2. CPD principles of operation 17 1.4.3. CPD measurements for charge detection: general considerations 19 1.4.4. The electrolyte-semi conductor interface 22 1.4.5. Modeling of CPD measurements of electrolytesemiconductor systems 24 1.5. Summary and dissertation organization 26 Chapter 2. SiC Surface Prepar ation and Characterization 29 2.1. H-etching of SiC surfaces 30 2.2. 3C-SiC 31 2.2.1. Effect of H-etching on the morphology of 3C-SiC surfaces and development of an ‘optimum’ etching process 32 2.2.2. H-etching rates of 3C-SiC(001) 40 2.2.3. Crystallographic studies and ch emical analysis of the near surface region: LEED, AES 41 2.3. 4H/6H SiC 46 2.3.1. H-etching processes fo r hexagonal SiC polytypes 47 2.3.2. Surface morphology: AFM 50 2.4. Summary 54 Chapter 3. SiC Biocom patibility Studies 56 3.1. Cell culture on 3C-, 4H -, 6H-SiC surfaces 57 3.1.1. Sample characteristics and cleaning 58


ii 3.1.2. Cell culture and expe rimental procedure 59 3.1.3. SiC superior biocompatib ility: MTT and fluorescent microscopy results 62 3.1.4. SiC vs. Si: evaluation of cell protrusions via AFM and optical microscopy 69 3.2. Influence of surface properties on cell adhesion and proliferation 71 3.2.1. Surface chemistry and wettability as possible explanations of SiC greater biocompatibility 72 3.2.2. Influence of SiC surface topography on cell adhesion and proliferation 75 3.2.3. Influence of SiC surface chemistry on cell adhesion and proliferation 77 3.3. Cleaning of SiC surfaces for bioapplications: RCA vs. Piranha 82 3.3.1. Effect of RCA and Piranha cleans on semiconductor surface morphology and chemistry 83 3.3.2. RCA clean as a promising surface pre-treatment for the study of cell-SiC adhesion sites 85 3.3.3. Effect of repeated Pira nha cleanings on chemistry, wettability and cell proliferation 88 3.4. Summary 91 Chapter 4. CPD Apparatus and Characterization of Surfaces for Cell-Semiconductor Electroni c Interaction Studies 93 4.1. CPD apparatus, calibration and experimental procedure for CPD measurements in air 94 4.1.1. Experimental apparatus 95 4.1.2. CPD system calibration and measurement precautions 97 4.1.3. Procedure for CPD measurements of semiconductors in air ambient 101 4.2. SiC and Si for cell-semiconducto r interaction studies: sample selection and description 102 4.2.1. Sample requirements for dry and wet CPD measurements 103 4.2.2. Selected SiC and Si samples for cell-semiconductor CPD investigations 104 4.3. Surface potential of SiC and Si in ‘steady state’ 105 4.4. Effect of hydrogen etching on the surface potential of SiC surfaces 108 4.4.1. Electronic passivation by Hetching of n-type 3C-SiC epilayers 109 4.4.2. Characterization of passivated 3C-SiC surfaces via XPS and ATR-FTIR 113 4.4.3. Surface potential of H-etched 6H-SiC 119 4.5. Effect of chemical treatmen ts on SiC and Si substrates 120 4.5.1. Chemical charging expe rimental procedure 120 4.5.2. Band bending operated by chemical charging of the surface: results and discussion 121


iii 4.6. Summary 124 Chapter 5. CPD Studies of the Semi conductor-Cell-Electrolyte System 125 5.1. Experimental procedure for CPD measurements of the semiconductor-cell-electrolyte system 127 5.1.1. Chemical preparation 127 5.1.2. Sample processing 128 5.1.3. Cell deposition / culture 131 5.1.4. Experimental procedure for CPD measurements of semiconductor-cell-electrolyte systems 133 5.2. Electrolyte semiconductor systems 134 5.3. Cell line selection and properties 139 5.4. CPD investigations of adherent mammalian cells-SiC systems 142 5.5. CPD investigations of RBC-SiC systems 147 5.6. Discussion and modeling of the obtained results 152 5.7. Summary 160 Chapter 6. Conclusion and Future Work 163 6.1. Conclusion 163 6.2. Future work 165 References 169 About the Author End Page


iv List of Tables Table 2.1. Etch rates of 3C-SiC(001) at different temperatures as resolved by FTIR measurement. 40 Table 2.2. AES data for the different structures observed on 3C-SiC(001). 44 Table 3.1. Surface roughness of semiconductor surfaces used in the biocompatibility study. 59 Table 3.2. Wettability of SiC and Si su rfaces as measured via sessile drop method. 74 Table 3.3. Elemental concentrations and ratios for 4H-SiC(0001) and 4H-SiC(000 1). 81 Table 3.4. Piranha effect on surface wettability assessed by sessile drop method. 89 Table 3.5. Si and C elemental concen tration and C/Si ratio for 4H-SiC samples treated with Piranha zero and ten times. 90 Table 4.1 CPD voltage mean ( cpdV) and standard deviation ( ) measured for the grounded chuck in the dark and at different probe-to-sample distances. 98 Table 4.2 Comparison of CPD measurem ents of 4H-SiC epilayers using two CPD tools. 101 Table 4.3 Semiconductor surface potentia l versus the overall surface charge. 106 Table 4.4 Charging effect of HF dip on H-etched n-type 3C-SiC epilayers. 113 Table 4.5 Effect of HF dip on Si and SiC surface potential as measured via CPD. 122 Table 4.6 Effect of potassium perma nganate on Si and SiC surface potential as measured via CPD. 122


v Table 4.7 Summary of the effect of HF and KMnO4 on Si and SiC surfaces: charge added by the chemical trea tment with respect to the initial state and sign of the surface potential measured. 123 Table 5.1 Effect of water on the surface potential of Si(111) and 3C-SiC(001) as measured via CPD. 136 Table 5.2 Effect of McCoy culturing media on surface potential of SiC and Si surfaces as measured via CPD. 137 Table 5.3 Effect of phosphate buffe r saline (PBS) on surface potential of SiC and Si surfaces as measured via CPD. 138 Table 5.4 Effect of different amount s of McCoy culturing media on surface potential of 3C-SiC(001) as measured via CPD. 138 Table 5.5 Non-detected electronic in teraction between HaCaT cells and 3C-SiC(001) after 24 hours from seeding (standard deviation < 15 mV). 145 Table 5.6 Non-detected electronic in teraction between B16-F10 cells and n-type 3C-SiC, n-type 6H-SiC and p-type 6H-SiC after 4 hours from seeding (standard deviation < 15 mV). 146


vi List of Figures Figure 1.1. Tetrahedron building block of all SiC crystals showing the bond lengths between Si and C atoms. 6 Figure 1.2. Illustration of the three different positions that the hexagonal frame of SiC bilayers can assume in the lattice (top) and stacking sequence of the three most common polytypes (bottom). 6 Figure 1.3. Energy band diagram for an ideal n-type semiconductor. 15 Figure 1.4. Energy band diagram of a real n-type semiconductor which contains an amount of band bending at the surface. 16 Figure 1.5. Schematic representation of a Monroe probe-grounded sample system. 18 Figure 1.6. Energy band diagram of an n-type semiconductor (a) not charged and (b) with the negative charges on the surface. 21 Figure 1.7. Electron energy levels for an electrolyte with respect to the vacuum level. 23 Figure 1.8. Electrical doubl e layer model describing the charge distribution at an n-type semiconducto r/electrolyte interface. 23 Figure 1.9. Equivalent circuit at the semiconductor/electrolyte interface where CH, CSS and CSC are the Helmholtz, surface state and space-charge capacitance, respectively. 25 Figure 2.1. 1010m AFM micrograph showing the typical surface morphology of 3C-SiC epilayer before H-etching treatment. 33 Figure 2.2. 55m AFM micrograph a nd 22m (inset) of as-grown 3C-SiC(001) epilayers showing atomic steps. 34 Figure 2.3. 55m AFM micrograph and 22m inset of a typical 3C-SiC surface after H-etching at 1200 C. 36 Figure 2.4. 22m AFM micrographs of the morphology of 3C-SiC surfaces etched at 1350 C (a) and 1300 C (b), respectively. 37


vii Figure 2.5. Plan-view SEM micrographs of typical defects observed on as-grown (a) and H-etched samples at 1200 (b), 1250 (c) and 1300 C (d), respectively. 39 Figure 2.6. Etch rates of 3C-SiC(001) ve rsus H-etch temperature as measured by FTIR. 41 Figure 2.7. UHV analysis of hydrogen etched 3C-SiC(001) samples. 43 Figure 2.8. AFM micrographs of 4H -SiC(0001) before H-etching. 50 Figure 2.9. AFM micrograph and extracted line profile showing the presence of atomic steps on the surface of H-etched 4H-SiC(0001) surfaces. 51 Figure 2.10. AFM micrograph of the surface morphology of on-axis 4H-SiC(0001) after H-etching. 52 Figure 2.11. Morphology and line profile of H-etched off-axis 6H-SiC(0001) surfaces. 53 Figure 2.12. AFM micrograph and line prof ile showing the atomically flat on-axis 6H-SiC(0001) surfaces obtained after H-etching. 54 Figure 3.1. AFM micrographs of (a) Si, (b) 3C-SiC grown on (100)Si, (c) 4H-SiC and (d) 6H-SiC surfaces. 58 Figure 3.2. Schematic representation of the sample positioning in the multi-well plate for cell plating. 61 Figure 3.3. Cell prolifera tion of B16, BJ and HaCaT cells expressed as x m measured via MTT assays at the third day. 63 Figure 3.4. HaCaT cell proliferation on Si and SiC substrates measured via MTT at the first and eighth day of culture expressed as x m. 64 Figure 3.5. Morphology of B16 ((a), (d)), BJ ((b), (e)) and HaCaT ((c), (f)) cells at the third day of cu lture on SiC and Si substrates, respectively. 65 Figure 3.6. Higher magnification images of the morphology of B16 ((a), (d)), BJ ((b), (e)) and HaCaT ((c), (f)) cells at the third day of culture on SiC and Si substrates, respectively. 66


viii Figure 3.7. Cell morphology of BJ fibroblasts at the fifth day of culture on a (a) SiC and (b) Si substrates as measured with flourescence microscopy. 67 Figure 3.8. B16 cell prolif eration measured via M TT at the second day of culture on 6H-SiC and GaAs substrates and expressed as x m. 68 Figure 3.9. Morphology of healthy B 16 cells cultured on 6H-SiC (a) vs. cytostructural degeneration of B16 cells cultured on GaAs as measured with flourescence microscopy. 68 Figure 3.10. Optical (left) and AFM (ri ght) micrographs of (a-d) B16 cells on SiC and (e, f) Si surfaces. 71 Figure 3.11. XPS surveys of two of the Si and SiC samples used in § 3.1.3. 73 Figure 3.12. Measured water contact a ngles on Si, 3C-, 4H-, and 6H-SiC surfaces. 74 Figure 3.13. AFM micrographs showi ng the morphology of the 3C-SiC samples used to evaluate the effect of surface roughness on cell adhesion and proliferation. 76 Figure 3.14. Patterned fibrobl asts adhesion on the sliver ed edge of a 3C-SiC sample which displays macroterraces 2000 to 4000 nm wide. 77 Figure 3.15. Comparison in HaCaT cell proliferation between etched and un-etched 3C-SiC substrates at the third day of culture expressed as x m measured via MTT assay. 79 Figure 3.16. HaCaT cell proliferation at the third day on Siand C-face 4H-SiC expressed as x m measured via MTT assay. 80 Figure 3.17. Optical images showing the surface morphology of 4H-SiC(0001) after cell growth and subsequent cleaning with (a) Piranha and (b) RCA cleans. 84 Figure 3.18. Presence of nitrogen (N) and sodium (Na) in the XPS survey of 3C-SiC(0001) after cell growth and subsequent RCA clean. 85 Figure 3.19. Analysis of RCA cleaned surface. 86 Figure 3.20. Analysis of different ad hesion sites found on SiC surfaces after RCA clean. 87


ix Figure 3.21. (a) Optical image and (b) AFM micrograph, with related line profile, showing a fibrillar ne twork on the SiC surface after RCA cleaning. 88 Figure 3.22. HaCaT cell proliferation at the third day on 3Cand 4H-SiC samples never cleaned with Pi ranha and cleaned with Piranha 10 times expressed as x m measured via MTT assay. 89 Figure 4.1. Experimental CPD appara tus housed inside a Faraday cage (probe, sample, chuck and LED shown). 96 Figure 4.2. Comparison between the Vcpd signal detected by the CPD system and the original voltage signal as observed via oscilloscope. 99 Figure 4.3. Probe charging effect on the surface of a 3C-SiC epilayer within the first 18 hours. 100 Figure 4.4. AFM micrographs (22 m scans taken in tapping mode) reporting the morphologies of the samples selected for CPD measurements: (a) n-type Si(111) (b) 3C-SiC(001), (c) n-type 6H-SiC(0001). 105 Figure 4.5. Band diagram representation of the ‘Steady state’ condition for nand p-type (a) Si and (b) SiC surfaces, respectively. 107 Figure 4.6. Surface potential vs. time of several H-etched 3C-SiC(001) epilayers presenting similar characteristics. 110 Figure 4.7. Surface potential vs. time of H-etched 3C-SiC(001) epilayers with final hydrogen cooling temperatures of 400 C (unfilled squares), 550 C (filled diamonds ), 1000 C (filled squares) and 1200 C (filled triangles). 111 Figure 4.8. Surface potential time monito ring, via CPD, of two Ar-annealed 3C-SiC(001) samples with similar characteristics. 112 Figure 4.9. XPS spectrum of a H-etched 3C-SiC(001) epilayer and relative elemental concentrations as calculated from the survey. 114 Figure 4.10. Si2p and C1s core level spectra obtained for the same 3C-SiC epilayer before etching (bold line) and after etching (light line). 115 Figure 4.11. ATR-FTIR spectra of H-etch ed 3C-SiC in a C-H stretch region indicating the existe nce of different typologies of C-H bonds. 117


x Figure 4.12. ATR-FTIR spectra of un-etch ed, HF treated 3C-SiC in a Si-OH stretch region displaying the existence of Si-OH bonds. 118 Figure 4.13. Surface potential time monitori ng, via CPD, of H-etched n-type (triangles) and p-type (c ircles) 6H-SiC surfaces. 119 Figure 5.1. Sample configuration for cell-CPD experiments showing (a) the free-standing sample approach and (b) the PEEK approach. 130 Figure 5.2. Cell morphology on (a) a sa mple mounted within the PEEK sample holder, (b) a free-standi ng sample in the vicinity of an epoxy drop (b). 131 Figure 5.3. Flow chart describing the experimental procedure adopted to investigate the effect of cell charge on semiconductor band bending. 144 Figure 5.4. Measured CPD surface potential for increasing amounts of RBCs for different 3C-SiC(001) samples. 148 Figure 5.5. Measured CPD surface potenti al for increasing amounts of RBCs for two n-type 6H-SiC(0001) samp les and data repeatability in two different experiments (DAY 1 and DAY 2) as calculated for one sample. 149 Figure 5.6. Measured CPD surface potenti al for increasing amounts of RBCs for two p-type 6H-SiC(0001) samples. 150 Figure 5.7. Repeatability of the surface potential values calculated for two different 3C-SiC(001) samples from two different experiments (DAY 1 and DAY2). 151 Figure 5.8. Schematic illustration of th e electronic status at the n-type SiC/electrolyte interface and re lative potential variation in the fluid. 153 Figure 5.9. Energy band diagram for n-type (LHS) and p-type (RHS) 6H-SiC/electrolyte interface assuming that the Fermi level position in the semiconductor is 200 meV from the conduction and the valence band edges, respectively. 154 Figure 5.10. Schematic representation of a negatively charged cell suspended in liquid and the system relative potential diagram. 155 Figure 5.11. Schematic illustrating the lim ited electronic effect of the cell charge on the semiconductor. 156


xi Figure 5.12. Schematic representation of the scattering and absorption of photons by the hemoglobin contained in the cell. 160 Figure 6.1. XPS spectra of 3C-SiC samp les with and without cell adhesion proteins. 167


xii Silicon Carbide Biocompatibility, Surface Control and Electronic Cellular Interaction for Biosensing Applications Camilla Coletti ABSTRACT Cell-semiconductor hybrid systems are a potential centerpiece in the scenery of biotechnological applications The selection and study of promising crystalline semiconductor materials for bio-sensing applicati ons is at the basis of the development of such hybrid systems. In this work we introduce crystalline SiC as an extremely appealing material for bio-applications. For the first time we report biocompatibility studies of different SiC polytypes whose results document the biocompatibility of this material and its capability of directly interfacing cells without the need of surface functionalization. Since the successful implementation of bi osensors requires a good understanding and versatile control of the semi conductor surface properties, th e chemistry, crystallography and electronic status of different SiC surfaces are extensively studied while their surface morphologies are thoroughly cont rolled via hydrogen etching. Also, investigations of the effect of cell surface charge on the electron ic status of SiC su rfaces are attempted adopting a contactless surface pot ential monitoring technique. The results obtained from these contactless measurements lead to the de velopment of theoretical models well-suited for the description of cell-semiconductor hybrid systems electr onic interactions.


1 Chapter 1. Introduction 1.1. Research objective and motivation Cell-semiconductor hybrid systems represent an emerging topic of research in the biotechnological area with intriguing possible a pplications. A comprehensive understanding of the interactions governing su ch systems is the basis of present and future development of biologically interfaced device performance. To date, very little is known about the main processes that govern the communication between cells and the surfaces they adhere to. When cells adhere to an external surface an eterophilic binding is generated between the cell adhesion protei ns and the surface molecules. After they adhere, the interface between them and th e substrate becomes a dynamic environment where surface chemistry, topology, and electron ic properties have been shown to play important roles [1-3]. Although previous works have demonstrated that cells display a net charge on their external surf ace [4-7] little is known about cell-semiconductor electronic interactions. Studying how and in which magnit ude the electronic properties of biological entities such as cells may influence and interact with the electronic status of a crystal, and describing it with the means provided by solid -state physics would be an enormous step forward in science and w ould represent the foundation for successful future implementation of electrically based bio-sensors. For this purpose, a suitable crystalline material displaying both biomedical and sensing potentialities should be selected and its properties fully characterized. A direct


2 interfacing of cells with the selected semiconductor is a re quirement for the detection, study and modeling of electronic signals. Howeve r, to date, the bioc ompatibility of only a few crystalline semiconductors has been investigated, with Si and titanium dioxide (TiO2) drawing most of the attention [3, 8-12]. In f act, since the present trend in research is the functionalization and polymer coating of semiconducting surfaces, simple biocompatibility studies of crystalline se miconductor materials have been mostly neglected. Both Si and Ti are unsuitable fo r the purpose of studying cellular electronic interactions with semiconductors. Si has b een shown to display different degrees of cytotoxicity, mostly due to its instability in aqueous soluti ons with subsequent formation of silica and silicates, which are known for their harmful effects on cells [9, 13, 14]. On the other hand, TiO2, which can become a semiconducto r upon ion implantation [3], does not display sufficiently satisfying electronic pr operties that may justify its adoption for electronic sensing applications. Therefore, it appears evident th at there is a need for the introduction of a different semiconducting mate rial that, displaying both biocompatibility and great sensing potentiality may fill the existing gap. Single-crystal silicon carbid e (SiC) is a wide band ga p semiconductor with vast sensing potentiality, very resistant to wear and corrosion, and with optimal tribological properties. In the past, because of its chem ical inertness and superior resistance, the amorphous phase (a-SiC) of this promising ma terial has been suggested for prosthesis and implant coating [15, 16]. For the same reasons, a-SiC biocompatibility has been widely studied yielding promising results [17-19]. Also, a-SiC ha s been found to be highly haemocompatible and an optimal coating for heart stents [20]. Although amorphous SiC has been widely characterized by bio-medical resear ch, surprisingly, to


3 date, no studies report on the biocompatibility of crystalline SiC, which, because of its wide electronic energy band gap, appears particularly appealing for bio-sensing applications. For the reasons listed above, we selected crystalline SiC as the ideal substrate material for bio-sensing investigations which may uncover the complex nature of semiconductor-cell electronic communication. Sin ce the degree of success in the use of a semiconductor material for biosensing applications str ongly depends on its biocompatibility and surface properties, we exhaustively studied crystalline SiC biocompatibility and characterized SiC surf aces at a chemical, crystallographic and morphological level. The interesting and insightf ul results obtained in the course of these studies are reported in the next chapters a nd, for the most part, are novel contributions that range from surface scien ce to biomedical fields. Cell-semiconductor electronic interaction in vestigations were al so part of this project. For these studies, we used fully char acterized SiC surfaces in combination with a contactless surface potential monitoring techni que: contact potential difference (CPD) measurements. The CPD technique was select ed because of its extremely appealing capability of monitoring the poten tial of a surface without di scharging it. The main idea at the basis of this experiment was to mon itor the effect that the cell charge has on the electronic status (e.g., energy band bending surface potential) of a semiconductor. Specifically, its implementation consists of monitoring, via CPD, the surface potential variation caused by the presence of cells culture d or deposited directly on SiC substrates. The success of such measurements would, as mentioned before, greatly impact the modeling and design of el ectronic biosensors. However, to th e extent that the final goal is


4 attractive, its implementation is equally challenging. One of the major obstacles involved with the CPD monitoring of cell-semiconductor sy stems resides in the f act that cells need to be immersed in liquid (e.g., culturing medi a) to be kept alive. This requirement introduces a significant challenge since the fe w CPD measurement attempts in the past using objects immersed in liquid have provi ded conflicting results [21-23]. Another issue related with the design of these CPD experiments is the very limited knowledge that we have presently on the effective charge of a cell immers ed in an electrolyte. The only experimental technique currently used with success which is capable of detecting the cell surface charge is electrophoresis [5, 24, 25] However, because of theoretical and modeling problems, the electrophor etic data do not a llow a direct calcul ation of the cell charge. As a result, except for erythrocytes, no estimations of the surface charge of cells are found in the existent literature. The lack of a defin ite quantification of this charge adds difficulty to the proper design of CPD experiments. Thanks to a suitable experimental approach we were able to overcome electrolyterelated experimental issues and to successf ully perform CPD measurements on surfaces immersed in liquids. This allowed us to electronically characterize the semiconductorelectrolyte interface and move towards the desired cell-semiconductor electronic interaction investigations. Our attempts to monitor possible changes in the semiconductor band bending due to the presence of live cells are also reported in this work and provide interesting information and the basis for fu ture cell-semiconductor electronic interaction studies. In this chapter we first describe the general characteristics of single-crystal SiC and the present knowledge of the bio-medical pot entialities of its amorphous phase (§ 1.2).


5 We then give a brief overview of all techni ques used in our work to characterize SiC surfaces (§ 1.3) and focus in particular on the CPD monitoring technique which is used later in this work for cell-semiconductor elec tronic interact ion investigations (§ 1.4). 1.2. Silicon carbide: a promising materi al for bio-sensing applications To the best of our knowledge, this is the fi rst work that introduces crystalline SiC as a promising biomaterial for bio-sensing applic ations. In the following sections we review the basic properties of this interesti ng semiconductor (§ 1.2.1) and the background information related to the use of its amorphous form, a-SiC, in biomedical research (§ 1.2.2). 1.2.1. SiC general properties Naturally occurring silicon carbide (SiC), which has the gem name of moissianite, was first observed in 1893 by Henri Moissan in the Canyon Diablo meteorite in Arizona. SiC is, in fact, extremely rare in nature and typically found only in minute quantities in corundum deposits, kimberlite, and meteorites Analysis of SiC gr ains contained in meteorites revealed that this material originated from stars evolved in higher-metallicity regions than our galaxy and suggested that Si C is older than our solar system [26]. To date, more than 170 crystalline forms of Si C can be synthesized, which is a property called polytypism. A tetrahedron of four car bon atoms covalently bonded to a silicon atom in the centre is at the basis of ever y SiC crystal (Fig. 1.1). Each carbon atom is located 3.08 from the others, while the di stance between the silic on and the carbon atom is 1.89 .


6 Figure 1.1. Tetrahedron building block of all SiC crystals showing the bond lengths between Si and C atoms. An interesting feature of SiC is that it displays a two-dime nsional polymorphism called polytypism. All the SiC crystals have a hexagonal frame of Si and C bilayers. The stacking order between succeeding double layers of carbon and silicon atoms is a variable that defines the different pol ytypes of SiC. Speci fically, in Fig. 1.2 th e three different positions that the hexagonal frame can assume in the lattice are reported and referenced as A, B, and C. Figure 1.2. Illustration of the three different positions that the hexagonal frame of SiC bilayers can assume in the lattice (top) a nd stacking sequence of the three most common polytypes (bottom).


7 What is defined as the cubic polytype ( or 3C-SiC) in SiC presents a stacking sequence ABCABC.... The hexagonal polytypes, 4Hand 6H-SiC ( -SiC), have stacking sequences of ABCBABCB… and ABCACBABC ACB…, respectively (Fig. 1.2). Due to the different crystallographic structure, different polytypes present different electronic and optical characteristics. The electron en ergy band gap is 2.39 eV for 3C-SiC, 3.023 eV for 6H-SiC, and 3.265 eV for 4H-SiC [27]. The properties that make this material particularly promising for biosensing applications are: 1) the wide bandgap that, as mentioned before, increases the sensing capabilities of a semiconductor; 2) the chem ical inertness that suggests the material resistance to corrosion in harsh environments such as body fluids (e.g. SiC does not react with any known material at room temperatur e, the only efficient etch being molten KOH at 400-600 C); 3) the high ha rdness (5.8 GPa), high elastic modulus (424 GPa), and low friction coefficient (0.17) that make it an ideal material for smart-implants and in-vivo biosensors [15, 28, 29]. As we mentioned above, the three ma jor polytypes of SiC present different properties. For completeness in this work we characterized the surfaces and evaluated the biocompatibility of the most studied SiC polytype (e.g., 3C-, 4H-, 6H-SiC). 1.2.2. SiC as a biomaterial: background information Even though crystalline SiC biocompatibility has not been investigated in the past, information exists concerning the biocompa tibility of the amorphous phase of this material (a-SiC). First, it is important to mention that a-SiC is one of the principal materials of choice for cardiovascular applica tions. In fact, a-SiC haemocompatibility has


8 been not only suggested by numerous studi es [29-31] but also greatly proved by the successful use of a-SiC coated heart stents in in-vivo clinical trials. In fact, to date, hydrogen rich amorphous SiC (a-SiC:H, also know n as Tenax) coated stents have been implanted in thousands of moderate-to-high risk patients yielding only a minor incidence of adverse events [20, 32-35]. At pres ent, a-SiC:H is known for its high thromboresistance induced by the optimal barrie r that this material presents for protein (and therefore platelet) adhesion. Also, amorphous SiC superior tribological and mechanical properties, togeth er with the fact that it exhibits hydroxyapatite-like osseointegration, make it an excellent candidate material for medical prosthetic implants [16, 36]. Although in-vivo trials of SiC coated hip or oral prostheses have not been performed (or at least reported) to date, in-vitro preliminary studies have shown promising results [16, 17, 37]. On the other ha nd, one of the possible drawbacks that may be associated with the use of SiC in-vivo is related to the unc lear and highly debated cytotoxic level of SiC partic les [38, 39, 17, 18]. Nonetheless, we believe that the potential cytotoxicity of SiC particles does not represent a dramatic issue as much as it does for Si, since the great tribological properties of SiC make it less likely to generate debris. In conclusion, most of the studies conduc ted in the past on a-SiC provide evidence of the attractive bio-potentialities of this ma terial and hence suggest similar properties for crystalline SiC. The promising information found in the literature combined with the well-known sensing potentiality of single-crystal SiC are at the basis of our choice to investigate the potentialities of this material for bio-sensing applications.


9 1.3. Surface characterization tools We previously mentioned that a successful implementation of a material for biosensing applications largely depends on the chemical, crystallogra phic and morphological properties of its surface. In f act, it is present knowledge th at the bonding of cell adhesion proteins to the molecules of a foreign surf ace is mostly influenced by the chemistry, hydrophilicity, morphology, and el ectrostatics of the surface itself [1-3]. For these reasons, this work also focused on the characterization of the surface of the semiconducting material that we propose as id eal for bio-sensing ap plications, namely SiC. In order to study SiC surfaces, we used different characterization techniques whose main features are reported below: Atomic force microscopy (AFM) operates by measuring the atomic forces between a sharp probing tip (typically a few nm in ra dius) and a sample surface. Even though it can be updated for different uses (e.g., magne tic characterization in magnetic force microscopy (MFM), doping profiling in sca nning spreading resistance microscopy (SSRM), etc.) the basic AFM system is us ed for morphological characterization of surfaces. The AFM can operate in several mode s: contact (the cant ilever contacts the surface while experiencing repulsive van der Wa als forces), non contact (the cantilever is held above the surface and senses the attract ive van der Waals forces), and tapping mode (e.g., the cantilever vibrates at or near its re sonant frequency and ‘taps’ the surface). The AFM basic principle of operation is the followi ng: a cantilever with a sharp tip is dragged across the sample surface while a laser is focu sed on the back side of the cantilever. The vertical probe motion is then sensed by a positionsensitive photodetect or and a feedback loop adjusts the probe-sample separation to ma intain a constant amplitude and force on


10 the probe. Hence, the feedback loop gives a m easurement of the sample height variation. The measurements are typically performed in atmosphere but, with proper equipment, can also be performed in vacuum. All the AFM results presented in this work were obtained working under atmosphere and, if not otherwise specified, in tapping mode. For the imaging in liquid, the contact mode was used. Scanning electron microscopy (SEM) is a technique for qualitative morphological characterization of a sample surface. An el ectron beam with elect ron energies ranging from 10 to 30 keV is raster-scanned across th e sample surface and the resulting electrons emitted from the sample are collected to form an image of the surface. SEM resolution is lower than that obtainable from AFM. Therefore, for high resolution and quantitative information the AFM method is preferred. SEM is typically used in this work to obtain large-scale characterization of sample surf aces followed by quantitative evaluation via AFM. Low energy electron diffraction (LEED) is used to investigate the crystallography of sample surfaces. Specifically, low energy (10-1000 eV) electrons incident on the sample are diffracted by the periodic arrange ment of the atoms. The diffracted electrons emerge from the surface in directions satisfyi ng interference conditions from the crystal periodicity and strike a fluor escent screen, forming a distin ct array of diffraction spots due to the orientation of the crystal lattice of the sample [40]. Obviously, because of their low energy, the incident electrons penetrate onl y a few atomic layers into the surface. The LEED results presented in Chapter 2 were obtained in an ultra-high vacuum (UHV) analysis chamber at the Max Planck Institut e (Stuttgart, DE) equipped with LEED optics,


11 an electron spectrometer for AES, a Si eva porator and sample h eating facilities. The pressure in the chamber during the measurements was in the 10-11 mbar range. Auger electron spectroscopy (AES) is a chemical characterization technique based on the Auger effect which describes the phe nomenon in physics for which the emission of an electron from an atom causes the filli ng of a vacancy in an inner electron shell. Specifically, primary electrons with typically 1-5 keV energy are emitted from an electron gun and impinge on the studied surface. If its energy is sufficient, an incident electron can remove a core state electron from a surface atom. This now empty core state can be filled by an outer shell electron from the same atom, in which case the electron moves to a lower energy state. The energy as sociated with this transition is typically imparted to a second outer sh ell electron, the Auger electron, which hence is ejected from the atom. The characteristic energy of this ejected electron defines the originating atom. By analyzing the resulting energy spectra it is possible to determine the chemical composition of the studied surface with th e exception of hydrogen and helium, which, having less than three elec trons, cannot be detected by this technique. The sampling depth of AES ranges from 0.5 to 5 nm [40]. Th e AES results presented in Chapter 2 were obtained in the same ultra-high vacuum (U HV) analysis chamber used for the LEED investigations. The energy of the incident electron beam was 2.25 keV. The pressure in the chamber during the measurements was again in the 10-11 mbar range. X-ray photoelectron spectroscopy (XPS) is a chemical characterization technique based on the photoelectric effect which allows as AES, identification of elements and their chemical status, with the exception of hydrogen and helium (in theory hydrogen and helium could be detected by using a very good spectrometer). Pr imary X-rays with


12 energies of 1 to 2 keV (e.g., Mg K radiation or Al K radiation) impinge on the sample surface causing ejection of electrons from a ny orbital. Obviously, photoemission occurs only for X-ray energies exceeding the elec tron binding energy. Therefore, each element generates a characteristic set of peaks in th e photoelectron spectrum at kinetic energies determined by the photon energy and the resp ective binding energies. Analysis of the peaks allows determination of the compositi on of the sample surf ace. Like AES, XPS allows sampling depths in the 0.5-5 nm range. All the XPS results presented in this work were obtained working at a maximum pressure of 10-9 mbar. Total attenuated reflectance Fourier transform infrared spectroscopy (ATR-FTIR) is a chemical characterization technique wh ich uses IR to determine the molecular composition of a surface. The FTIR principle of operation is base d on the fact that part of the incident infrared radiation is absorbed by molecules only if the frequency of the radiation provides energy in the exact quan tity required by the bond s to vibrate. The infrared spectra collected after the beam has passed through the sample is indicative of the chemical bonds present at the surface (i.e., different molecules absorb at different characteristic frequencies). Unlike AES and XPS, this technique is cap able of detecting H and its compounds. Attenuated total reflectance (A TR) is used in particular for obtaining IR spectra of difficult samples such as th e solid/liquid interface. ATR occurs when the incident beam enters from a more-dense (with a higher refractive index) into a less-dense (with a lower refractive index) medium. The evanescent wave from the primary optical beam penetrates a very short distance be yond the interface and into the less-dense medium before the complete reflection occurs (typically a few m). The wave intensity is reduced by the sample in regions of the IR spectrum where molecular absorption takes


13 place. To utilize the ATR effect, the sample is placed in close contact with a highrefractive-index optical crystal. A 45 bevele d zinc selenide (ZnSe) crystal was used in the ATR-FTIR experiments performed in this work. Atmospheric CO2 and H2O absorption lines were reduced by purging th e apparatus with dry nitrogen [41]. Other surface characteriza tion techniques and experi mental apparatus were sporadically used during the course of this research project, and will be introduced as the corresponding result s are presented. 1.4. Contact potential difference techni que for cell-semiconductor electronic interaction studies Contact potential difference (CPD) measuremen t is the technique selected in this research project for investigation of cell-se miconductor electronic interactions. In this section we first introduce semiconductor en ergy band diagrams, which are of primary importance for a complete understanding of the CPD technique (§ 1.4.1), and then discuss the principle of ope ration behind CPD measuremen ts (§ 1.4.2). In § 1.4.3, we describe how CPD measurements can be us ed to detect charges on a semiconducting surface. Also, since the CPD investigations we intend to perform involve a liquid layer which wets the semiconductor surface, it is important to discuss energy band diagrams for liquids and the existent theory for semi conductor/electrolyte interfaces (§ 1.4.4) and to use this information to model the CPD measurements of semiconductor-electrolyte systems (§ 1.4.5).


14 1.4.1. Semiconductor energy band diagrams The energy spectrum of an ideal semiconducto r presents two different typologies of energy levels: the allowed energy levels, situ ated within the conduction and the valence bands, and the forbidden energy levels, situated within th e so-called band gap. Energy band diagrams of semiconducto rs are charac terized by: the conduction band edge, Ec, which is the lowest unfilled energy level in the conduction band; the valence band edge, Ev, which is the uppermost filled energy level in the valenc e band; the Fermi level, EF, which represents the maximum energy occupied by an electron at 0 K; the energy band gap, Eg, which is given by Ec – Ev and contains all the forbi dden energy levels; and the vacuum energy level, E0, which is a reference level re presenting the energy of a free electron. The basic band diagram for a semi conductor is sketched in Fig. 1.3. Besides those already listed, different quantities appear in this figure: the difference between E0 and EF is called the work function, q and represents the average energy required to extract an electron from the semiconducti ng surface; while the difference between E0 and Ec is called the electron affinit y, qX, and is the energy necessar y to free an electron from the conduction band edge. In an intrinsic se miconductor (where electrons and holes are present in equal amounts), the Fermi level is near the middle of the forbidden gap and is called the intrinsic Fermi level (Ei). However, for n-type se miconductors (i.e., electrons are majority carriers) EF > Ei and for a p-type semiconductors (i.e., holes are majority carriers) EF < Ei. This basic information will help the reader to understand better the energy band diagrams presented in the remainder of this work.


15 Figure 1.3. Energy band diagram for an ideal n-type semiconductor. Fig. 1.3 represents the energy band diagram of an ideal semiconductor, where it is assumed that the allowed energy states at the surface are not different from those in the bulk. However this assumption is not valid for a real semiconductor, where the asymmetric nature of the crystal at the su rface (atoms at the surface are only single-side bonded) and the existence of crystal defect s and foreign bonded atoms introduce extraallowed energy states. These allowed energy states are known as surface states and vary in energy and typology. For example, states th at are neutral when occupied by electrons and positively charged when unoccupied are cl assified as donor stat es. States that are negative when occupied but neutral when empty are classified as acceptor states. This nomenclature will be often used in the followi ng chapters. The major effect that surface states have on an energy band diagram, and therefore on the elec tronic status of a semiconductor, is that of generating a ba nd bending (e.g., the surface is naturally charged). In fact it is the change in the band bending near semiconductor surfaces that is typically exploited in sensing applications a nd thus the key physical concept to take away from this introductory discussion.


16 Now that we have presented the basics of semiconductor energy band diagrams, we introduce the concept of surface potential that is the quantit y measured by the dark/light CPD technique to be presented later. The surface potential, s, is a measure of the semiconductor surface departure from the state of electrical neutrality, and is measured as the energy difference between the conduction band (valence band) ed ge at the surface and the conduction band (valence band) edge in the undisturbed part (e.g., the bulk) of the semiconductor. From this definition it is clear that an ideal semiconductor presents a null surface potential, while the surface poten tial of a real semiconductor is not zero because of the existence of surface states. In fact surface states cause a natural charging of the semiconductor surface with subsequent depletion or accumulation of majority carriers within the surface region. The band diag ram for a real n-type semiconductor, and its relative surface potential, is reported in Fig. 1.4. Figure 1.4. Energy band diagram of a real n-type semiconductor which contains an amount of band bending at the surface.The dotted lines (small dots) represent the ideal condition for Ec/q, Ev/q and i. The surface potential s is indicated. As is evident from Fig. 1.4, the surface poten tial can equally be defined using as a reference level the intrinsic potential, i. The surface region where depletion or accumulation of the majority carriers take place is named the space charge region (SCR).


17 The electronic condition where the SCR is depleted of majority carriers is called depletion and corresponds to an upward bending of the bands for n-type ( s < 0, Fig. 1.4) and a downward bending of the bands ( s > 0) for p-type semiconductors. In the case where the majority carriers accumulate at the surface (e.g., the condition named accumulation), the bands bend down ( s > 0) for n-type and up ( s < 0) for p-type semiconductors. The condition for which the intrinsic potential and the surface potential coincide is called flatband and obviously in this case s = 0. The treatment of these concepts will be resumed later in Chapter 4. 1.4.2. CPD principles of operation Surface potentials of a semiconductor ar e commonly measured via dark/light contact potential difference (CPD) measurements. In a CPD apparatus, a pick-up electrode (i.e., probe) is placed near the semiconductor surface, whose back-side is grounded, hence forming a capacitor, C, with the semiconductor. Either a Monroe or Kelvin probe can be used for this purpose [ 42]. Since in the appa ratus implemented for this work we used a Monroe probe [43], which typically pr esents a low sensitivity to external vibrations, we now describe th e Monroe probe-CPD pr inciple of operation. The final value obtained from a CPD me asurement is the contact potential difference, Vcpd, between the probe and the sample. Th is voltage can be easily detected by a Monroe probe-electrostatic voltmeter co mbination in the fashion we now describe. In the Monroe probe, the electrode is fixed and a grounded shutter, mounted in front of the electrode, is vibrated horizontally, varyi ng the area of the capacitor plates and thereby modulating the probe-to-wafer capacit ance as shown in Fig. 1.5 (e.g., C = A/d where C


18 is the capacitance, A is the ar ea of each plate, d is the se paration between the plates and is the permittivity of the insulator between the plates). Figure 1.5. Schematic representation of a Monroe probe-grounded sample system. Because of the vibrating shutters the cap acitance is time variable. Since I = Vcpd dC/dt, an alternating current (ac) is generated in the el ectrode. The contact potential difference Vcpd is determined by the electronics of the electrostatic voltmeter by applying the null-arrangement, initially proposed by Kelvin [44] Specifically, the current generated in the probe is nullified by adjusting the bias voltage VB until I = 0, in which case VB = Vcpd. Summarizing, a CPD apparatus determines the contact potential difference that exists between the probing el ectrode and the investigated semiconductor by means of a current-nulling method applie d on the current signal generated by a variable capacitor. Let us now introduce the prin ciple of operation of dark /light CPD measurements and explain how this technique yields the su rface potential of the analyzed sample. In dark/light CPD two contact potential difference values are measured: Vcpd,dark which is the value measured as described before in dark, and Vcpd,light which is the value obtained


19 under deep illumination of the sample. Assumi ng that the chuck-semiconductor contact is ohmic the CPD value measured in dark is a sum of different c ontributing potentials: Vcpd,dark = VFB + Vair + s (1) where the flatband voltage, VFB, is a function of the meta l-semiconductor work function difference and Vair is the voltage drop associated with the air. Moreover, if an oxide is present at the semiconductor surface the cont act potential difference measured at the probe is given by: Vcpd,dark = VFB + Vair + Vox + s (2) where Vox is the voltage drop associated with the oxide. As is evident from (1) and (2), a single CPD measurement is insufficient to determine the surf ace potential of a semiconductor. This is why measurements under deep illumination are performed. LightCPD measurements are performed using a li ght source with a phot on energy higher than the energy band gap of the semiconductor under study. Therefore, el ectrons are excited from the valence band to the conduction band (i.e., electron hole pair (EHP) generation) and force the semiconductor to a flatband condition ( s 0). Hence the contact potential difference measured by the probe under intense illumination is: Vcpd,light = VFB + Vair + Vox (3) Therefore the surface potentia l can be easily calculated as the difference between these two measurements: s = Vcpd,dark – Vcpd,light (4) 1.4.3. CPD measurements for charge de tection: general considerations It is now clear what a surface potential is and how dark/light CPD measurements detect it. Presently, this technique is widely used for the determination of the effect of


20 chemical treatments and ionic charges on a semiconductor surface. In fact, charges deposited or trapped on the semiconductor su rface induce a bending of the energy bands. The entity of this band bending, and therefore the sign and magnitude of the charge at the surface, can be determined by performing CPD measurements of the surface not charged and comparing the measured s value with the one obtained for the same charged surface. Let us now consider an n-type semiconduc tor in a naturally depleted condition (e.g., majority carrier absence in the SCR) as the one depicted in Fig. 1.6(a). By adding negative charges on its surface we will enhanc e the magnitude of the band bending (e.g., majority carriers are pushed deeper into the semiconductor bulk because of repulsion from the negative charge at the surface, Fig. 1.6(b)). CPD measurements of the surface in these two different conditions wi ll yield different values of s. Supposing that we did not know the sign of the deposited charge, a shift of s towards more negative values upon charge addition would have been a direct indication that the charge deposited was negative. Upon previous corona charge char acterization of the specific semiconducting surface we would also be able, by comparing the two measured s values, to define the magnitude of the deposited charge [45].


21 Figure 1.6. Energy band diagram of an n-t ype semiconductor (a) not charged and (b) with the negative charges on the surface. Note the difference in magnitude of the two svalues. Only one thing could actually impede an estimation of the amount of charge deposited on the surface: surf ace state density. In fact, wh en measuring the surface potential via CPD we are actually measuring the capacitor formed between the probe and the semiconductor surface. On the semi conductor surface (e.g., one plate of the capacitor), the charge is a sum of the charge in the SCR and that in the surface states. Therefore the capacitance measur ed via CPD is a parallel combination of the surface state capacitance Css and the space charge capacitance Csc. The presence of a high density of surface states may cause Fermi level pinning [1 3, 46], in which case charge addition on the semiconducting surface will not result in a response of excess charge in the SCR and the value of Css will be predominant in the capacitance measured via CPD. As we already pointed out our goal is to implement the CPD technique for detecting the charge associated with cells once th ey are cultured/deposited on SiC surfaces. Obviously, this can be achieved only if the surfaces we use presen t a reduced amount of surface states. This issue will be tr eated later in Chapters 2 and 4.


22 1.4.4. The electrolyte-semiconductor interface As explained in the previous section, CP D dark/light measurements of a bare semiconductor surface allow determination of the semiconductor surface potential. However, in this project we aim to perf orm CPD measurements of a more complex structure, the semiconductor-cel l-electrolyte system, for the determination of the surface potential variation induced in a semiconductor by the presence of ch arged cells. For this purpose we will perform CP D measurements of the se miconductor-electrolyte and semiconductor-cell-electrolyte systems and co mpare the surface potentials calculated for the two cases. Eventual differences in these valu es would allow us to define the effect of the charge of cells on the elec tronic status of a semiconducto r. To be able to understand the results obtained from CPD measurements of semiconductors immersed in electrolytes we first need to model the semiconductor/el ectrolyte interface. For this purpose we introduce an existent model th at electronically characterize s this complex interface and subsequently apply it for the modeling of our CPD measurements (§ 1.4.4). Let us first introduce the distribution of en ergy levels within an electrolyte as we did for a semiconductor in § 1.4.1. An electrolyte is electronically characterized by: a redox energy level, ERedOx, which defines the average ener gy level at equilibrium of all the individual redox species; the most proba ble energy level of the reduced species, ERed; and the most probable energy le vel of the oxidized species, EOx. A schematic representation of the energy band diagram for an electrolyte is re ported in Fig. 1.7 and, also in this case, referenced to the energy vacuum level, E0.


23 Figure 1.7. Electron energy levels for an electro lyte with respect to the vacuum level. Immersion of a semiconductor in an electroly te results in a charge transfer process (i.e., electron exchange) between the two phases until equilibrium is obtained, that is when the Fermi level in the semiconductor and the redox level in the electrolyte are equal (EF = ERedOx). This produces an electric field at the semiconductor/electrolyte interface which generates an electrical double layer we ll described by the Stern model depicted in Fig. 1.8. Figure 1.8. Electrical double laye r model describing the charge distribution at an n-type semiconductor/electrolyte interface. On the electrolyte side the position of the cl osest approach of mobile ions is called the outer Helmholtz plane (OHP). The Helm holtz layer is the region between the


24 semiconductor surface and the OHP and contains ions attracted to the semiconductor surface by the excess charge in the space ch arge region and by polar water molecules. Outside of the Helmholtz layer a region with excess ions of one sign whose thickness depends on the electrolyte concentration exists the so-called Gouy layer [13, 47]. As is evident from Fig. 1.9, the Gouy layer is a diffuse space charge region with excess ions of the same sign as those accumulating in the OHP. 1.4.5. Modeling of CPD measurements of electrolyte-semiconductor systems Now that we have introduced the existe nt model for semiconductor/electrolyte interfaces we can apply this knowledge to mode l the CPD measurements in liquid that we will describe in Chapter 5 and that may lead to a better understanding of the semiconductor-cell elec tronic interaction. When measuring, via CPD, a semiconductorelectrolyte system different voltage drop contributions will sum to yield the Vc pd value measured by the Monroe probe. To ensure that the value obtained from the differe nce of dark and light measurements is the surface potential (i.e., s = Vcpd,dark Vcpd,light) all the additional vol tage drops should be constant values independent of the illumination. Furthermore, we must define their magnitude to ensure that their contributi on is not predominant in the measured Vcpd value. An equivalent circuit for the electrical components at the semiconductor/electrolyte inte rface is reported in Fig. 1.9.


25 Figure 1.9. Equivalent circuit at the semi conductor/electrolyte interface where CH, CSSand CSC are the Helmholtz, surface state and sp ace-charge capacitance, respectively. The voltage drop between the semiconducto r bulk and its surface is the value s, which for a dry semiconductor can be dire ctly calculated via dark/light CPD. As discussed in § 1.4.3, this voltage drops acros s two capacitances in parallel, the space charge capacitance, Csc, and the surface states capacitance, Css. In a semiconductorelectrolyte system s will also be in series with the voltage drops associated with air (Vair), with the Helmholtz layer (VH) and with the Gouy layer (VG). However, for electrolyte concentrated solutions (e.g., above 10-2 M), the value of VG is negligible [48]. Therefore the final electronic eq uivalent circuit in terms of the relevant capacitances will be given by the series of the Helmholz capacitance CH with the parallel of the space charge capacitance Csc and the surface states capacitance Css, as shown in Fig. 1.9 and described by equation (5). sc ss H ss sc H cpdC C C ) C C ( C C (5) Therefore it is evident that in order to monito r the effect that the ch arge associated with cells has on the electronic status of a se miconductor (e.g., changes to its space charge region) we need: Css<

26 solution like the ones used in this work are close to 20 F/cm2 [49] and typical values of Csc for SiC samples range from 10-2 to 10-1 F/cm2 (depletion to accumulation, respectively) the latter condition can be c onsidered satisfied. Also, the value of Csc is influenced by the doping density in the SiC cr ystals and, therefore, can be controlled. However Css, which is determined by the surface st ate density, is another matter and is highly dependent on how the crystals are prep ared and their surfaces treated. Therefore reduction of surface state density on SiC su rfaces was attempted using hydrogen etching (Chapter 2) in order to ensure that the above condition was met. 1.5. Summary and disse rtation organization The electronic interactions that exist between biological cells and semiconductors are, to date, unknown. In this work we aim to move towards a be tter understanding of how, and in which magnitude, the charge of a cell may influen ce the electronic status of a semiconductor. In order to accomplish this very ambitious objective a semiconductor that combines both biocompatibility and sensing pote ntialities needs to be selected and fully characterized. In this chapter we introduced crystalline SiC as a promising material for this task: its wide band gap is extremely appealing for sens ing applications and the wellknown biocompatibility of amor phous SiC likely suggests that the crystalline phase may display similar properties. In order to char acterize SiC surfaces different techniques were used in the course of this wo rk, which have been introduced in this chapter. Moreover we suggested an apparently promising technique to investigate cell -semiconductor electronic interactions: contact potential difference (CPD ) measurements. Since, to date, no studies report on the direct investigation of the e ffect of cell charges on semiconductors, the implementation of successful measurements that may lead to uncover the electronic


27 communication ongoing between the biolog ical and the semiconductor world is pioneering and extremely challenging. Also, although the CPD techni que appears to be ideal for our goal because of its contactless nature, its implementation for hybrid system monitoring is not trivial. In particular, the n ecessary presence of an electrolyte during the measurements complicates the matter and accu rate modeling of both the hybrid system, which is the object of measurements, and of th e measurement itself due to such artifacts, are required to better desi gn and understand the inves tigations we performed. The organization of the dissert ation, based on the individual chapters, is as follows. In Chapter 2 we introduce hydrogen etching as an ideal technique for obtaining high quality, atomically flat SiC surfaces with resu lting low surface state charge densities. The H-etched surfaces are also perfectly suitab le (e.g., well prepared) for surface science studies. This allowed us to characterize in a comprehensive fa shion the morphological, chemical, and crystallographic features of SiC surfaces by using AFM, AES, and LEED methods, respectively. Chapter 3 reports an exhaustive study on the biocompatibility of SiC. Fluorescence microscopy, viability assays and atomic for ce microscopy were used to characterize the morphology, adhesion quality and prolifera tion of mammalian cells on SiC substrates. This study, besides offering, for the first time, quantitative and qualitative information on the biocompatibility of SiC, also describes the possibility of directly interfacing cells with SiC surfaces. This result definitely confirms th e potentialities of Si C as a biomaterial and an ideal substrate for cell-semiconduc tor electronic in teraction studies. Chapter 4 describes the implementation of a CPD apparatus for cell-semiconductor electronic interaction studies and reports on the electroni c characterization of SiC


28 surfaces. The effect of charges on SiC substr ates is exhaustively investigated and the possibility of passivating elec tronically the surface of 3C-SiC is presen ted and discussed. Chapter 5 reports the procedures developed for CPD measurements of semiconductors immersed in liquids and investig ates the effect that different electrolytes have on the electronic status of SiC surf aces. Also, in this chapter we report the procedure and results for CPD measurements of semiconductor-cell-el ectrolyte systems. Chapter 6 summarizes the results reported in the previous chapters and uses the knowledge acquired in the course of this research to sugg est future developments and possible implementations of cell-semic onductor electronic inte raction studies.


29 Chapter 2. SiC Surface Pr eparation and Characterization In order to succeed in the implementation of sensors for biomedical applications the surfaces of the semicoducting materials that ar e going to be interfaced with cells have to fulfill several requirements. First, they need to be fully characterized at a chemical, crystallographic, and morphological level. It is in fact well known that cell adhesion to a surface is regulated by a combination of chemical, morphological and energetic properties of the surface [1-3]. Also, atom ically flat surfaces may be preferable, depending on the targeted application. Spec ifically, for the investigation of cellsemiconductor electronic interactions that we aim to perform, working with atomically flat surfaces is particularly appealing: th e reduction of surface defects and hence of associated surface states surely simplifie s the task by reducing trapping at the cell/semiconductor interface. In this section we present the processes developed for producing atomically flat SiC surfaces, and their chemical and crystall ographic characterization. Commercially available SiC crystals present slicing and po lishing scratches at the surface as a result of the wafer preparation processes, which imply a high density of surface states in addition to a high surface roughness. Moreover, these su rfaces are not feasible for chemical and crystallographic studies, which would allo w for their accurate characterization. A comprehensive understanding of the chemical structural and mo rphological surface characteristics is of primary importance for all those biosensing applications which, as


30 the CPD measurements proposed later in this work, rely on the direct interfacing of biological cells and semiconducti ng material. Hence, the need for a full characterization of the surfaces that will be used in the rest of this work becomes clear. Scratch-free, passivated SiC surfaces suit able for surface science studies can be obtained by using an appropriate surface preparat ion technique. The chemical inertness of SiC makes it impossible to use wet chemical etchants to remove the slicing and polishing damage. To date, the SiC surface preparation techniques which have been developed and used with different degrees of success are: oxidation [50, 51], sub limation etching [52], photoelectrochemical etchi ng [53], chemomechanical polishing (CMP) [54], plasma etching [55], and hydrogen etching (H-etching) [56-58]. Among these, the latter has been shown to be the most effective to completely remove polishing scratches while producing atomically flat surfaces pe rfectly suitable for surf ace science studies [58-61]. This section focuses on the production a nd characterization of well-ordered 3C-, 4Hand 6H-SiC surfaces via H-etching in a hot-wall chemical vapor deposition (CVD) reactor. Morphological, crystallo graphic and chemical analyses are performed via atomic force microscopy (AFM), low energy electr on diffraction (LEED) and auger electron spectroscopy (AES), respectively. 2.1. H-etching of SiC surfaces H-etching of SiC surfaces has been primarily used in the past to improve the surface quality of bulk substrates prior to epitaxial growth [62-64]. It is usually performed in a chemical vapor deposition (CVD) reactor (either hot-wall or cold-wal l) flowing variable hydrogen fluxes at high temperatures (typica lly in the range of 1000-1700 C), and at a pressure varying between 100-760 Torr.


31 Hydrogen etching of SiC is believed to be a two-step process [59, 63], which can be simplified as follows. The high temperature characteristics of the process cause the evaporation of Si atoms from the surface a nd the subsequent exposure of the underlying C atoms. At this stage, hydrogen atoms fr om the etching gas bond to the surface C atoms forming hydrocarbons which in turn desorb fr om the crystal uncovering the next layer of Si which then evaporates, and so on. Howeve r, the removal of ma ny monolayers (up to several hundred nanometers) of SiC material not always pr oduces atomically flat and ordered surfaces. First, the morphology of th e surface obtained after etching is highly dependent on the original surface condition. If the surface has a hi gh density of hidden defects or presents heavy polishing dama ge, the etching will expose and enlarge the subsurface defects and worsen the surface morphology. Also, one of the most common problems in SiC H-etching is the condensatio n of Si droplets on the surface due to the preferential hydrocarbon evaporation caused by the higher equilibrium pressure of C-H groups with respect to Si [62]. Also, over-e tching of the SiC surfaces, which results in step-bunching and defect enlargement, must be prevented to achieve an atomically-flat surface. In the following sections, we will use the terminology ‘optimum process’ to designate an etching process which impr oves the surface morphology and the atomic order of the starting surface without causing the aforementioned drawbacks. Careful planning of the etching experiment, together with an understanding of hydrogen etching kinetics allowed us to develop ‘optimum’ Hetching processes for 3C-, 4Hand 6H-SiC. 2.2. 3C-SiC In the past, the lack of sufficiently large hi gh-quality 3C-SiC substrates has led to a delay in the development of growth techniques for this polytype and a subsequent lack of


32 good quality 3C-SiC epilayers. As a conseq uence, the chemical, morphological and crystallographic characteristics of 3C-SiC epila yers have been investigated on a smaller scale than those of the SiC hexagonal polyt ypes (i.e., 4H/6H-SiC de scribed in § 2.3). Since no reports on H-etching of 3C-SiC su rfaces were found in the literature, we had no references to help us define the optimal etching parameters. Therefore, in order to develop the ‘optimum’ etching process, we fi rst studied the effect of H-etching on the morphology of 3C-SiC epilayers for different et ching parameters (§ 2.2.1). Etching rates of 3C-SiC(001) for different etching te mperatures are reported in § 2.2.2. Crystallographic studies of the 3C-SiC surf aces etched under optimum conditions showed a surface reconstruction (e .g., (51)) which has not been investigated previously in the literature and was therefore thoroughly exam ined in § 2.2.3. Also, the ‘optimum’ Hetching process presented in § 2.2.1 may be helpful in the development of a H-etching process for the removal of damage and scra tches on polished 3C-SiC surfaces after CMP treatment. 2.2.1. Effect of H-etching on the morphol ogy of 3C-SiC surfaces and development of an ‘optimum’ etching process The 3C-SiC(001) epilayers us ed in this work were gr own on 8x10 mm Si(001) dice in a low-pressure, hot-wall, horizontal ch emical vapor deposition (CVD) reactor [65] using a chlorinated growth chemistry [66]. Th e film crystallinity was confirmed by x-ray diffraction (XRD) rocking curves [66]. Noncontact doping profiling measurements were also performed which assessed the film doping to be n-type and in the low 1015 cm-3 range [67]. The thickness of the studied epilayers vari ed between 2 and 6 m, as measured by Fourier transform infrared spec troscopy (FTIR). Only the epilayers which


33 showed, during AFM inspection, similar morphological characteristics were used in this work. AFM micrographs representative of these as-grown 3C-SiC(001) surfaces are shown in Figures 2.1 and 2.2. Figure 2.1. 1010m AFM micrograph showi ng the typical surface morphology of 3CSiC epilayer before H-etching treatment. Note the presence of APDs. AFM data taken in tapping mode. Figure 2.1 reveals that the as-grown surfaces presented anti-phase domain boundaries (APDs), which are typi cal features of 3C-SiC f ilms heteroepitaxially grown on Si and are due to the large lattice mism atch (~20%) between the 3C-SiC and the Si crystals.


34 Figure 2.2. 55m AFM micr ograph and 22m (inset) of as-grown 3C-SiC(001) epilayers showing atomic steps. Step hei ght and terrace width along the [110] directio n are depicted in the higher resolution image a nd extracted by line profile. AFM data taken in tapping mode. As is evident from the 55 m micrograph in Fig. 2.2, the samples exhibited the atomic structure of the SiC crystal even immediately after growth. Within individual APDs the AFM images revealed atomically fl at terraces with small steps. However, the steps were only loosely ali gned along the low-Miller i ndex directions [110] and [ 110] and were rather wavy. Nevertheless, typical st ep heights, as determined from the higher resolution micrograph and the line profile also displayed in Fi g. 2.2, were in the 2-3 regime which corresponds well to biatomic step s (2.18 ) in the 3C-SiC crystal structure which in the <001> direction is characte rized by alternating C and Si layers. For reference the trenches observed between di fferent APDs were at least 15 nm deep. Four parameters can be varied in a H-etchi ng process: temperature (T), pressure (p), hydrogen flow (Hf), and etching time (tetch). In the set of expe riments that will be described in this section, which has been designed with the double aim of studying the effect of H-etching on cubic SiC surfaces a nd developing an ‘optimum’ etching process,


35 only temperature was varied from experiment to experiment, while the other parameters were kept constant. Since the epilayers to be etched presented APDs, aggressive etching had to be avoided in order to prevent prefer ential etching and subsequent enlargement of the trenches (i.e. sites with higher surface en ergy are subjected to higher etch rates). Low pressure (LP) etching is known to increase the etch rate (Si diffuses better through the thinner boundary layer caused by lower pre ssures): therefore, the whole set of experiments was performed at atmospheric pressure (AP) while Pd-purified hydrogen was flown at the moderate rate of 10 SLM (s tandard liters per minute). H-etching of a crystal operates in an opposite fashion than growth, therefore etching temperatures are typically comparable to those used for growth The 3C-SiC epilayers studied in this work were grown on Si(001) (whose melting temperat ure is 1410 C) at temperatures around 1385 C. However, for our experiments, we opted for a lower range of etching temperatures (between 1200 a nd 1350 C) since a low proce ssing temperature has fewer requirements on the reactor, lower cost and re duced contamination. An etching time of 30 min was chosen as a constant in all the experiments performed. To summarize, all the etching experiment s were performed at AP, flowing 10 SLM of Pd-purified hydrogen, for 30 min and with temperatures ranging between 1200 and 1350 C in a hot-wall CVD reactor specific ally designed for 3C-SiC growth and processing [68]. The reactor hot-zone graphite and parts used were not exposed to growth in order to minimize the possi bility that particles rem oved during the growth would deposit on the SiC surface. In order to rem ove the native oxide present on the 3C-SiC epilayers, the samples were dipped in a 50:1 mixture of de-ionized (DI) water and hydrofluoridic acid (HF), rins ed with DI water, and drie d with nitrogen. The samples


36 were then immediately loaded into the r eactor and brought under vacuum so as to minimize the growth of any native oxide on the crystal surface. AFM micrographs representative of the surfaces obtained after etching at 1200, 1300 and 1350 C are reported in Figures 2.3 and 2.4. Figure 2.3. 55m AFM micrograph and 22m inset of a typical 3C-SiC surface afte r H-etching at 1200 C. Step height and terrace width along the [110] direction are depicte d in the higher resolution image and extracted by line profile. AFM data taken in tapping mode. As is immediately evident by comparison of the three AFM micrographs shown in Figures 2.3 and 2.4, the best surface mor phology was obtained after etching at 1200 C. The samples etched at this temperature, in fact, presented a well-defined cubic morphology with steps perfectly aligned along the [110] and [ 110] directions, as shown in Figure 2.3. The more clear alignment of the atomic steps along the low-Miller index directions represent an evident improvement of the processed surface with respect to the as-grown (un-etched) surface of Figure 2.2. Again the terraces are atomically flat, 50–100 nm wide, and typically separated by 2–3 steps as seen from the higher resolution image and corresponding line profile in Figure 2.3. Step bunching to multiples of the 2.18


37 value was only occasionally observed, while the trenches of the APDs did not appear to be enlarged or deepened when compared to the as-grown (un-etched) samples. Figure 2.4. 22m AFM micrographs of th e morphology of 3C-SiC surfaces etched a t 1350 C (a) and 1300 C (b), respectively. Note the highly damaged surface after etching at 1350 C. As a reference, Rq = 35 nm for image (a) while is 0.88 nm for image (b). The line profile along the [110] dir ection is extracted from fi gure (b) and depicts the step height and terrace width. AFM data taken in tapping mode. On the other hand, a significant enlargement of defects (generally in the form of square-shaped pits) was observed in samples etched at 1300 C. At this temperature, step bunching becomes evident, as indicated by the hi gher step heights (0.8-1.2 nm) and larger terrace widths (150-400 nm) depicted in Fig. 2. 4(b) and in the extracted line profile. At 1350 C, a drastic change of the surface mor phology into a highly mosaic cubic structure was observed (Fig. 2.4(a)). As a refere nce, the root mean square roughness (Rq) of a 22m AFM micrograph of samples etched at 1300 C was typically less than 1 nm and comparable to the roughness of samples either as-grown or H-etch ed at 1200 C. Instead, the highly damaged samples etched at 1350 C presented a Rq higher than 30 nm. SEM studies, while adding new informa tion in a more macroscopic scale, confirmed what was revealed by AFM analysis SEM micrographs of samples etched at a) b)


38 1200, 1250, 1300 and 1350 C are shown in Figure 2.5. For as-grown samples, SEM analysis revealed a significant presence of two defect types on the surface: SiC cluster hillocks and film cracks. A micrograph of th e surface crack feature is shown in Figure 2.5(a). In samples etched at 1200 C we found a surface pit density sim ilar to the original surface crack density. Since the pit defects (Fig. 2.5(b)) were never found on as-grown surfaces it is logical to assume that these are enlargements of the surface cracks caused by hydrogen etching. In support of this assumption are the SEM micrographs of samples etched at 1300 C presenting very large and deep pits extending down to the Si substrate (Fig. 2.5(c)). Fig. 2.5(d) shows the highly damaged mosaic structure, already observed via AFM, of surfaces etched at 1350 C, whic h is caused by the aggressive preferential etching of surface defects and grain boundaries that takes place at this temperature. It has to be mentioned that surfaces etched at 1250 C were also studied via AFM and SEM. While AFM micrographs revealed no evident morphological differences with respect to the samples etched at 1200 C, SEM analysis showed a higher density of pits and defects on these surfaces.


39 Figure 2.5. Plan-view SEM micrographs of t ypical defects observed on as-grown (a) an d H-etched samples at 1200 (b), 1300 (c) and 1350 C (d), respectively. A growth hillock is present in the center of the highly damaged mosaic surface morphology depicted in (d) for reference [69]. From the reported results we can conclude that the ‘optimum’ process parameters for the 3C-SiC(001) epilayers studied in this work are the following: p = 760 torr, T = 1200 C, Hf = 10 SLM, and tetch = 30 min. However, for a more complete study of the effect of H-etching on 3C-SiC( 001) we also performed etchi ng processes where either the pressure or the etching time were varied with respect to the ‘optimum’ values. An etching time of 10 to 20 min seemed to produce litt le or no differences on the macroscopic surface morphology. However, the resulting at omic steps appeared to be less clearly aligned with the [110] and [ 110] directions for shorter etching processes. Also, as expected, LP processes provoked a more aggr essive etching of the surface defects and APD trenches.


40 2.2.2. H-etching rates of 3C-SiC(001) The etching rate of 3C-SiC(001) was ev aluated at different temperatures by measuring the thickness of the epilayers befo re and after etching via Fourier Transform Infrared Spectroscopy (FTIR). For this expe riment the etching time was fixed at 1 hour, the pressure at 760 torr, Hf was 10 SLM while the etching temperature was varied between 1200 and 1375 C. From this set of experiments we obs erved an exponential dependence of the etching rate on temperature as shown in Fig. 2.6. The mean etch rate and standard deviations were calculated fr om three etching experiments conducted at each temperature and are reported in Table 2.1. Very small etch rates were observed when etching at temperatures below 1300 C. Also, it appear s that the ‘optimum’ etching process (1200 C for 30 min) was actually re moving just few monolayers of material. It has to be pointed out that while the thickness variation for samples etched at temperatures higher than 1300 C was easily resolved by FTIR measurement, the etch rate determination at lower temperatures was affected by the limited sensitivity of the FTIR, the relative flatness of the samples, and the difficulty to perform the measurement always in the same position on the sample. Table 2.1. Etch rates of 3C-SiC(001) at different temperatures as resolved by FTI R measurement [69]. Values reported as mean standard deviation. T (C) 1200 1250 1300 1325 1350 1375 Etch rate sd (nm/hr) 17.5 6.2 21.3 11.875.1 9.9170 2.8 344.5 25.9 580.5 45.9


41 Figure 2.6. Etch rates of 3C-SiC(001) versus H-etch temperature as measured by FTIR. Note exponential dependence of etch rate vs temperature with only a few monolayers o f material removed at the optimum temperature of 1200 C [69]. 2.2.3. Crystallographic studies and chemi cal analysis of the near surface region: LEED, AES The surface structure and the chemical co mposition of the near surface region of the 3C-SiC(001) samples etched with the ‘optimum’ process were i nvestigated in a ultra-high vacuum (UHV) analysis chamber at the Max-Planck-Institute (Stuttgart, DE). This was accomplished via low-energy electron diffraction (LEED) and Auger electron spectroscopy (AES) analysis conducted on a set of 4 samples. The UHV chamber was equipped with LEED optics, an electron spec trometer for AES, a Si evaporator and sample heating facilities. The pressure in the chamber during the measurements was in the 10-11 mbar range. The symmetry and location of the diffraction spots in the LEED pattern were used to reveal surface order and superstructure periodicities. Comparison of peak intensities of the differentiated AES spectra served to determine relative concentrations of elements in the near surf ace region. The samples were investigated in UHV in their native state after hydrogen etching and subsequen tly after Si deposition and


42 annealing at different temperatures. As-grown samples, cooled down under Ar flow, were also studied for comparison. After loading into the UHV chamber and w ithout any intermediate treatment, the samples revealed a sharp LEED pattern of ap proximate (51) periodicity. Fig. 2.7(a) shows the superposition of the diffraction pa tterns of two domain or ientations of this structure rotated by 90 with respect to ea ch other. A similar LEED pattern has been briefly noted by Kaplan [70] to occur during the partial reduction of a strongly oxidized 3C-SiC(001) surface, however without being further inte rpreted. In our study we observed that the long unit vector of the supercell was not ex actly an integer multiple of the substrate (11) unit vector. Rather, the superstructure was incommensurate and its unit vector varied on different samples of nominally equal prep aration. We found supercell sizes varying in the range from (4.51 ) to (6.51). It appeared that the substrate and the overlayer fail to have a precise ep itaxial relationship. For convenience, the periodicity of the structure is denoted as “5x1” from this point forward in the text. The chemical composition of the “51” structur e was investigated using the AES spectra shown in Fig. 2.7(c). Most notably, the as in troduced surface contained oxygen as seen in the bottom curve, indicating either the formati on of a thin oxide layer after air exposure or, similar to the case of the SiC(0001) and SiC(000 1) surfaces, silicate layer formation during the etching process [56] The significant intensities of the Si and C peaks suggest the oxide thickness to be limited to one or a few atomic layers at most. As inferred from the AES spectra in Fig. 2.7(c) (top curve), the oxygen was completely removed from the surface by flashing the sample in UHV to about 1120 C for 5 min. Concurrently, the “51” diffraction spots disappeared from the LEED pattern in favor of a poorly ordered


43 version of the known (21) reconstruction [71]. Evidently, the “51” phase represents an ultra-thin reconstructed oxidic layer on the 3C -SiC(001) surface. It should be noted that this “51” pattern was found on both H-etched and as-grown epi-layers, and even after RCA cleaning of an as-grown sample. Figure 2.7. UHV analysis of hydrogen etch ed 3C-SiC(001) samples. (a) LEED patter n of the “5x1” phase as obser ved after hydrogen etching (wit hout any further treatment). (b) LEED pattern of the (32) phase obtained after Si deposition a nd annealing to 1070 C in UHV. (c) AES spectra with the Si, C and O transitions indicated for the hydrogen etched and the annealed sample, and (d) for the (32), (21) and c(22) phases [72]. The composition of this oxide phase wa s determined from a comparison with established surface phases on 3C-SiC(001) pr epared by Si deposition and subsequent annealing. After stepwise (5 min each) an nealing of the Si enriched surface to temperatures in the range of 700-1100 C, the LEED pattern showed a (32) surface


44 reconstruction with sharp and intense 1/3 order spots on the integer rows, i.e. (m/3,n) spots, and streaky contributions on half order position, i.e. the (m/3,n+1/2) spots, the latter of which were best de veloped at about 1070 C (Fig. 2.7(b)). The AES spectra in this temperature range revealed that the su rface was still Si-rich (Fig. 2.7(d), bottom curve). At 1170 C the (32) transformed into a well-ordered (21) phase. The AES spectrum indicated a lower concentration of Si for this phase (Fig. 2.7(d), middle curve). Annealing to 1230 C provoked further preferen tial desorption of Si from the 3C-SiC surface and produced a C-rich surface (top curv e in Fig. 2.7(d)) with a c(22) LEED pattern. All three phases are well known on 3C-SiC(001) [71]. The (32) structure is believed to be strongly Si-rich [73], as also shown by our results. The (21) reconstruction is also reported as a Si-rich phase [74], while the c(22) phase is usually interpreted as a C-terminated structure [75, 76], which again is consistent with the AES development in our studies. Table 2.2. AES data for the different struct ures observed on 3C-SiC(001). Peak-to-pea k amplitudes evaluated for diffe rentiated Si, C and O AES signals and element intensity ratios. For comparison data are listed for the so-called silicate layer reconstructions on the basal plane surfaces SiC(0001) and SiC(000 1) taken from ref. [56]. H-etched “51” 1120 C diffuse (21) (32) (21) c(22) SiOx-Siface[56] SiOx-Cface[56] Si 1 0.9 1.2 0.6 0.7 1.4 1.7 C 1.3 0.9 0.3 0.5 0.75 0.75 1.4 O 0.35 0.225 0.10 Si/C 0.77 1.0 4.0 1.2 0.95 1.9 1.2 O/Si 0.35 0.16 0.06


45 Table 2.2 summarizes the peak-to-peak amp litudes of the differentiated Si, C and O AES signals and element intensity ratios for a ll structures observed here on 3C-SiC(001). In addition data are recalled for the so-calle d silicate layer reconstructions on the basal plane surfaces SiC(0001) and SiC(000 1) taken from ref. [56]. Consideration of the Si/C intensity ratio clearly places the “51” recons truction on the less Si rich side of the phase diagram. With some caution one might even favor a carbon termination underneath the oxidic layer. On the other hand the surface might contain some hydrocarbon contamination and the observed C enrichment may not be entirely due to the SiC structure. However, the data definitely rules out that the observed “51” structure is a derivative of the UHV prepared (52) phase whic h is known to be strong ly Si rich and is characterized by a Si-dimer rec onstruction layer [77]. The O/Si ratio observed is slightly higher than for the silicate reconstructions on the basal plane surfaces, which at a first glance suggests the presence of more than one monolayer of oxide. However, on the basal plane surfaces the Si/C ratio is larger than in our hydrogen etched case which, applying the above argument again, could suggest that, after correction for possible carbon contamination, the oxide thickness is si milar, i.e. one monolayer in our case. Thus, as a feasible scenario for the oxide related “51” reconstruction we suggest the presence of a nearly bulk terminated SiC surface with an oxygen containing reconstruction layer of the observed “51” pe riodicity. The varying interface periodicity suggests a relatively weak bonding relations hip between the substrate and the oxide overlayer, which in turn might be due to an (at least partially) internally passivated SiOx network with a certain energetic degeneracy for the distance of adjacent oxide-substrate bonds.


46 Finally, we should recall the presence of the two-domain superposition in the “51” LEED pattern. Due to th e alternating layer stacking se quence in the crystallographic <001> direction of 3C-SiC(001), the subs trate surface possesse s only a two-fold symmetry, which would necessarily imply that only one of the domains can be present. Provided that the interface does not form ar bitrarily on both Si and C surface terminating layers, the presence of both domains can on ly indicate that the 3C -SiC APD’s can have both possible epitaxial orientations with respect to the Si(001) substrate, [100] || [100] or [100] || [010]. This is an important observa tion for possible further optimization of the growth process. 2.3. 4H/6H SiC Unlike 3C-SiC, H-etching of the SiC he xagonal polytypes (4Hand 6H-SiC) has been largely studied and practiced during th e past couple of decades [56-59, 62-64, 78, 79]. Typical etching apparatus are tantalum strip heaters [80], vapor phase epitaxial (VPE) reactors [81], hot-wal l [62] and cold-wall CVD reactors [58] where the SiC etching temperatures range between 1400 and 1600 C. Low pressure etching or chlorinated chemistry (HCl addition) can be used to increase the etch rate [59]. However, HCl addition has been found to cause preferen tial etching at low etching temperatures (e.g. 1400 C) [59, 81], and in general highe r etching rates can wo rsen the resulting surface if the original substrate is highly sc ratched or defective. Because of the large amount of information which can be found in th e existent literature [56-59, 62-64, 80, 81] the design of the etching pro cesses for 4Hand 6H-SiC has been more straightforward than for 3C-SiC. This implies that no studies of the effect of temperature and pressure on the crystals was necessary in this work. In fact atomically flat surfaces were obtained


47 both for 4Hand 6H-SiC after the first etch ing attempt, which was probably due to an accurate choice of the etching pa rameters based on careful studi es of the data found in the existent literature and insightful considera tions concerning the kinetics in the CVD reactor used for the processing. 2.3.1. H-etching processe s for hexagonal SiC polytypes Hydrogen etching of both 4H and 6H-SiC samples was performed in a hot-wall CVD reactor mostly used for SiC epilayer growth [65]. The hydrogen flown during the process was Pd-purified. The substrates used in this study were: eight 4H-SiC(0001) 8 off-axis, of which half were chemo-mechan ical polished (CMP) by Novasic; two on-axis 4H-SiC(0001); two 3 off-axis 6H-SiC( 0001); two on-axis 6H -SiC(0001). All the samples were 810 mm diced from SiC bulk crystals purchase d from Cree Research, Inc, with n-type doping ranging from 10-18 to 10-19 cm-3. Also in this case the samples were dipped in H2O:HF (50:1) and rinsed in DI water just prior to the etching experiments. Etching of off-axis surfaces is known to be more problematic than for on-axis. Theoretically, one would expect a faster ma terial removal on off-axis surfaces where there is a higher density of high surface ener gy zones (e.g. steps edges, because of the narrower terraces). Instead, several studies re port higher etching rates for on-axis rather than for off-axis SiC surfaces and suggest the use of higher temperatures or more aggressive etching processes to completely remove the polishing scratches from off-axis surfaces [58, 59, 62, 82]. Also, the morphology of off-axis surfaces is reported to be less ordered than that of on-axis, with terraces much narrower (because of the larger off-axis angle) and steps not perfectly oriented along the low-index Miller directions [59, 62].


48 In this work, we first tried to develop tw o etching processes, one for off-axis 4HSiC and one for off-axis 6H-S iC, and successfully adopted them also for etching on-axis samples. Surprisingly, both the designed et ching processes produced independently from the presence of a miscut, atomically flat 4H and 6H-SiC surfaces, wh ich are described in the next section. Even though the on-axis surf aces present some tr iangular peninsulas between steps, which may be caused by the use of an etching process specifically designed for off-axis surfaces and therefore mo re aggressive, their surface quality after etching is still satisfactory. This suggests th at the presented etching processes are quite versatile for the etching of both on-axis and off-axis hexagonal SiC surfaces. Also, it has to be pointed out that all the samples us ed presented polishing scratches prior to Hetching and that those scratches were co mpletely removed by the adopted etching processes. The etching parameters for the off-axis processes were chosen in an attempt to minimize the etching temperature and based on the following considerations. Successful etching processes in hot-wall CVD reactors we re reported to operate at temperatures between 1500 and 1600 C for off-axis 4H-SiC and above 1600 C for off-axis 6H-SiC, with a pressure of 760 Torr, and an etchi ng time between 10 and 80 minutes [62, 64]. The higher etching temperatures needed for 6H-SiC may be explained by the different growth conditions adopted for the two polytypes. The CVD reactor used in our etching processes presents a peculiar tapering of the susceptor, which has been designed to avoid the preferential consumption of growth precursors at the susceptor entrance and to maintain a constant gas concentration throughout the w hole growth-zone by incr easing the velocity of the gases along the flow direction [83]. As a result, the higher gas flow generates, for


49 the same etching parameters, a more effective etching. Hence, we chose etching temperatures of 1400 C for 4H-SiC and of 1550 C for 6H-SiC, while the pressure was fixed at 760 Torr, the H-flow at 10 SL M, and the etching time at 30 min for both processes. As pointed out at the beginning of this se ction, the CVD reactor used for etching hexagonal SiC was mostly a ‘growth-dedicated’ reactor. In order to prevent the formation of the typical defects (e.g., Si-droplets, particulate deposition, undulate surface patterns) usually found in the crystal surface after etch ing with a ‘growth-dedicated’ CVD reactor, clean susceptor and poly-crystalline SiC samp le support plates (e .g., not exposed to growth) were used. However, the graphite foam used in these experiments was also used in the past for growth experiments, increasi ng the risk of particle contamination which can nucleate triangular surface defects during et ching. The presence of residual growth precursors in the reactor can in fact interfere with the initial stage of etching causing an initial localized growth at some areas on the substrate either due to the incomplete deposition of species or to lo calized etching of th e surface. Initial lo calized growth can then result in triangular (or ‘undulate’) patter ns on the surface as reported in [62]. For these reasons, the precaution of baking the CVD reactor in hydrogen with all its etching components was taken. The baking was performe d for 1 hour at atmospheric pressure and 1400 C before each 4H/6H etching experiment.


50 2.3.2. Surface morphology: AFM The 4H off-axis samples which underw ent the H-etching process presented different initial morphological features: ha lf, chemo-mechanical polished by Novasic, displayed during AFM analysis an extremely fl at surface but no atomic steps (Fig. 2.8(b)) while the remaining presented typica l polishing scratches (Fig. 2.8(a)). Figure 2.8. AFM micrographs of 4H-SiC(0001) before H-etching. (a ) typical polishing scratches on the ‘regularly’ polished surfaces; (b) flat surfaces of the chemo-mechanical polished samples. Note the smaller scale in (b) helps to visualize the absence of atomic steps. As a reference, the deepest scratc hes in (a) are 20 nm deep. AFM micrographs taken in tapping mode. After etching, both sets of samples presen ted atomically flat surfaces with terraces 50-100 nm in width and steps half a unit cell in height (5 ) (Fig. 2.9). The steps were aligned along the [1 100] direction while unit cell steps were seldom observed.


51 Figure 2.9. AFM micrograph and extracted lin e profile showing the presence of atomic steps on the surface of H-et ched 4H-SiC(0001) surfaces. Etch parameters were: T = 1400 C, Hf = 10 SLM, p = 760 Torr and tetch = 30 min. AFM micrographs taken in contact mode. On the other hand, the on-ax is 4H-SiC samples, which before etching presented a morphology like the one depicted in Fig. 2.8 (a), presented much wider atomic terraces after being etched with the same process used for the off-axis surfaces. This result is in total agreement with what has been reported in the literature [62]. The terraces observed for the on-axis surfaces were 1 m wide, while the steps had unit cell height (Fig. 2.10). Also in this case the steps were aligned along the [1 100] direction. In Fig. 2.10 it is evident the presence of triangular peninsulas protruding from the steps onto the terraces below them. Their height was assessed to be half a unit cell via AFM. These peninsulas are explained in [63] in terms of “fast-etch ” and “slow-etch” direct ions. We suggest that they may be an indication of a slight ove r-etching of the 4H-SiC on-axis surfaces.


52 Figure 2.10. AFM micrograph of the surf ace morphology of on-axis 4H-SiC(0001) afte r H-etching. Relative line profile defines the step height to be 1 nm and the terrace width 1 m. Note the presence of tr iangular peninsulas. Etch pa rameters were: T = 1400 C, Hf= 10 SLM, p = 760 Torr and tetch = 30 min. AFM micrographs taken in tapping mode. Similarly to the 4H-SiC surfaces, both the on-axis and off-axis 6H-SiC(0001) surfaces displayed a morphology as the one report ed in Fig. 2.8(a) before etching. After etching at 1550 C for 30 minutes both the o ff-axis and on-axis 6H-SiC(0001) presented atomic steps during AFM inspection. Also in this case, as for 4H-SiC, the steps were observed to be aligned along the [1 100] direction and the ones on off-axis surfaces were much narrower than the ones for on-axis crystals, as expected. A typical AFM micrograph of the atomically flat surfaces obt ained after etching off-axis 6H-SiC samples is reported in Fig. 2.11. The extracted line pr ofile indicates that th e steps are between 5 and 7.5 high and the terraces are 40-80 nm wi de. These values are comparable to those reported in [59, 64] for off-axis 6H-SiC su rfaces. Since one unit cell of 6H-SiC is 15.12 high in the direction perpendi cular to the (0001) plane, th e steps observed after etching were most certainly half a unit cell high (7.5 ). The 5 high steps observed in Fig. 2.11 are probably an underestimation of half a unit cell steps due to the difficulty to resolve with precision the height of narrow terraces with the software analysis program used.


53 Figure 2.11. Morphology and lin e profile of H-etched offaxis 6H-SiC(0001) surfaces. Etch parameters were: T = 1550 C, Hf = 10 SLM, p = 760 Torr and tetch = 30 min. AFM micrograph taken in contact mode. The typical morphology observed for H-etch ed on-axis 6H-SiC(0001) samples is reported in Fig. 2.12. As is evident, the st eps are half a unit cell high and terraces are 150 nm wide. The same values have been report ed by Owman [64] for H-etched on-axis 6HSiC(0001) surfaces. The two particles observe d in Fig. 2.12 are probably caused by the ‘growth-dedicated’ carbon insulating foam used during the etching pr ocess to support the clean graphite susceptor. It has to be men tioned that particulate is often found on the surface of samples etched with the CVD react or described in this section even after baking of the system. This finding strongly confirms the need of an ‘etch-process dedicated’ foam insert for the aforementioned CVD reactor.


54 Figure 2.12. AFM micrograph and line profile showing the atomically flat on-axis 6HSiC(0001) surfaces obtained after H-etching. The presence of particulate is due to the use of a ‘growth-dedicated’ foam in the CVD reactor. Etch parameters were: T = 1550 C, H f = 10 SLM, p = 760 Torr and tetch = 30 min. AFM microgra ph taken in tapping mode. 2.4. Summary As discussed in Chapter 1, SiC is a promising candidate for biotechnological applications. A successful implementation of th ese applications puts specific demands on the quality and surface propert ies of SiC. In particular the CPD measurements proposed later in this work, which may help uncove r cell-semiconductor elect ronic interactions, require well-ordered, chemically passivated and fully characterized surfaces. For this reason, this chapter reports on the control and understandin g of the major polytypes of SiC via H-etching and surface characterization techniques. H-etching is shown to be a successful technique able to produce atomically flat and re peatable SiC surfaces. Exciting possibilities such as functionalization and nano-patterning of SiC surfaces with biomolecules become more feasib le thanks to the morphologica l atomic order revealed by H-etching treatment. Also, for those applicatio ns which require a di rect interface between the biological cell and the semiconducting ma terial, as the CPD measurements described


55 in this work, a comprehensive chemical a nd crystallographic characterization of the semiconducting surface, as reported in § 2.2.3, becomes of primary importance. In conclusion, the preparation and complete ch aracterization of atomically ordered SiC surfaces may lead to the successful implementa tion of a large variety of biotechnological applications and, in particular, of the CPD studies proposed in this work and presented later in Chapters 4 and 5.


56 Chapter 3. SiC Biocompatibility Studies In Chapter 2 we have discussed how suitable surface prepar ation and accurate surface characterization are of fundamental importance for semiconducting materials used in bio-sensing applications. However, besides properties such as low surface roughness and surface state densities there is another requirement that a semiconductor that is going to interface cells needs to fulfill: biocompatibility. As we already mentioned in Chapter 1, despite the promising potentiality for biosensing applications, no studies are found in the literature which investigate crystalline SiC biocompatibility. In this chapter, we study single-crystal Si C biocompatibility by culturing mammalian cells directly on SiC substr ates and by evaluating the resulting cell adhesion quality and prolifera tion. We also compare SiC biocompatibility to the one of the leading crystalline semiconductor in biot echnology, Si, whose cytotoxicity has been reported by several studies [84-86]. Nonetheless, Si continues to be widely used because of the ease of electronic integration with biological systems (e.g., arrays for retina implant, micro resonator prob es for cortical recording, etc.). The results of this biocompatibility study show that crystalline SiC is indeed a very promising material for bio-applications, with better bio-performance th an crystalline Si. This result opens up exciting perspectives for the use of SiC for bi o-technological applicat ions. In particular 3C-SiC, which can be directly grown on Si substrates, appears to be an especially promising bio-material: the Si substrate used for the epi-growth woul d in fact allow for


57 cost-effective and straightforward electroni c integration, while the SiC surface would constitute a more biocompatib le and versatile interface between the electronic and biological world. Also, the excellent capabili ty of SiC in directly interfacing with biological cells without the need of a ny surface functionaliza tion is of primary importance for the successful implementation of the CPD measuremen ts proposed later in this work and an optimum starting point for the investigation of cell-semiconductor electronic interactions. Since a material’s biocompatibility is in fluenced by chemical, morphological and electrostatic fa ctors we also explore, in th is chapter, the effect that different surface chemistries, morphologies and wettabilities have on cell adhesion and proliferation. The importance of using an appropriate cleaning procedure for the SiC samples before their use as substrates for cell cultures is also discussed. 3.1. Cell culture on 3C-, 4H-, 6H-SiC surfaces A reliable and relatively easy way to de fine a material’s cytotoxic level, biostability, possible bioactivity and overall biocompatibility is by culturing cells directly on it while monitoring the cell adhesion quality and prolifera tion. In this section we evaluate the biocompatibility of different SiC polytypes by culturing three lines of mammalian cells directly on 3C-, 4H and 6H-SiC surfaces and by using MTT [3-(4,5dimethylthiazol-2-yl)-2,5-diphe nyltetrazolium bromide] assays and optical imaging to monitor cell proliferation and adhesion quality, respectively. In § 3.1.1 we describe sample characteristics and cleaning prior to cell seeding. Section 3.1.2 reports on the experimental procedure adopted for culturi ng cells and evaluating the material biocompatibility while § 3.1.3 discusses the MTT assay and fluorescent microscopy results. Additional studies investigating di fferences in cell adhesion morphology for Si


58 and SiC samples were performed via optical microscopy and atomic force microscopy (AFM) and are discussed in § 3.1.4. 3.1.1. Sample characteristics and cleaning The samples used as substrates for cell cu lture were as follows: three 3C-SiC(001) epilayers grown on Si(001) in the CVD reactor described in § 2.2; three 4H-SiC(0001) off-axis bulk crystals H-etched with the process reported in § 2.3.1 and originally purchased from Cree, Inc.; three 6H-SiC(0001) bulk crystals and Si (001) bulk crystals also purchased from Cree, Inc. All samples were diced to have a dimension of 8x10 mm. The doping was n-type for all the samples and in the range of: 1015 atoms/cm3 for 3CSiC; 1018 atoms/cm3 for 4H-SiC; and 1019 atoms/cm3 for 6H-SiC and Si. The surface characteristics of the 3C-SiC epilayers were those of the as-grown samples described in § 2.2.1. The 4Hand the 6H-SiC samples pres ented surfaces as described in § 2.3.1 for etched 4Hand un-etched 6H-SiC samples, respectively. The morphologies of these surfaces are depicted in Fig. 3.1 while their root mean square roughness (Rq), measured via AFM for an area of 25 m2, are listed in Table 3.1. Figure 3.1. AFM micrographs of (a) Si, (b ) 3C-SiC grown on ( 100)Si, (c) 4H-SiC an d (d) 6H-SiC surfaces. All the micrographs are 22 m in dimension. AFM data taken in tapping mode.


59 Table 3.1. Surface roughness of semiconducto r surfaces used in the biocompatibility study. Surface roughness estimated from AFM measurements taken in tapping mode. Substrate material Si(001) 3C-SiC(001) 4H -SiC(0001) 6H-SiC(0001) qR(55 m) 0.8 nm 1.4 nm 0.5 nm 1.2 nm As is evident from Table 3.1, all the samples presented similar roughness values (0.5 nm < Rq (25 m2) < 1.4 nm) and, therefore, eventual differences in measured cell adhesion quality and proliferation cannot be ex plained in terms of di fferences in surface topology. Some minor experime nts whose results are reported in the following sections also involved gallium arsenide (GaAs) and 4Hand 6H-SiC(000 1) (nominally C-face) samples. For reference, these samples presented surface roughness values comparable to the ones reported in Table 3.1 but since the majority of data presented here does not involve these samples they have not been included in the table to avoid confusion. All the samples which were repeatedly us ed as substrates for cell culture were cleaned of any organic residue by imme rsion in a piranha solution (2:1 H2SO4:H2O2) for 5 minutes. After a rinse in de-ionized (DI) water, the surface oxide generated by the immersion in Piranha was removed by dipp ing the samples in a hydrofluoric acid solution (50:1 H2O:HF) for 2 minutes. The samples were then thoroughly rinsed in DI water, dipped in ethanol, rinsed again a nd finally dried with dry nitrogen. All H2O was deionized (DI) with a resi stivity of at least 16 Mcm. 3.1.2. Cell culture and experimental procedure The mammalian cells used were as follo ws: 1) B16-F10 mouse melanoma cells (ATCC CRL6475); 2) BJ human fibrobl asts (ATCC CRL2522) and 3) human


60 keratinocytes cells (HaCaT). The B16 a nd BJ cells were purchased from ATCC. Different media were used in the culture of the different cell li nes: Eagle’s Minimum Essential Medium (EMEM) supplemented with 10% fetal bovine serum (FBS) for the BJ cells; McCoy’s Modified Medium supplemen ted with 10% FBS for the B16 cells; and Dulbecco’s Modified Eagle’s Medium (DME M) supplemented with 10% FBS for the HaCaT. Prior to cell plating on the semiconducto r surfaces, all the cells were cultured in 25 cm2 culture flasks (Corning), incubated at 37 C in an air atmosphere containing 5% CO2 and 95% relative humidity, and split and/or used at confluence. Immediately before cell plating, three samp les for each substrate material / polytype examined in the assay (e.g Si, 3C-, 4Ha nd 6H-SiC) were cleaned as described in the previous section and then plac ed in a 15.6 mm diameter cell culture well within a multiwell ultra-low cell-attachment plate (Corning). A schematic re presentation of the sample positioning in the multi-well plate is sketched in Fig. 3.2. Three empty wells, like the one represented in Fig.3.2. only belonging to a multi-well culturing plate (Corning), were used as controls for each MTT assay performed. It is evident that the plating areas are different for the semiconducting sa mples and the control wells. As a point of reference of the 200 mm2 well area less than 2/5 were covered by the 80 mm2 semiconductor samples. Therefore, the seeding density values were scaled accordingly for the plating area as described in [87]. Cells were seeded in th e wells with the semiconducting samples at the following densities: 40x103 cells/cm2 for the BJ cells; 15x103 cells/cm2 for the B16 cells; and 20x103 cells/cm2 for the HaCaT. Obviously, separa te experiments were conducted for the three different cell lines. After seeding, all the plated cells were incubated at 37 C for 1 to 8 days in an air atmosphere containing 5% CO2 and 95% relative humidity.


61 Figure 3.2. Schematic representation of the sample positioning in the multi-well plate for cell plating. A multi-well ultra-low cell-attachment plate was used during all experiments reported in this diss ertation with a well area of 200 mm2. Seeding was performed in triplicate to allowstatistical analysis of the resu lting cell viability an d adhesion. A separate multi-well plate was seeded with cells and used as control. In the schematic all the seeded wells are drawn within the same multi-well plate for ease o f representation. Cell proliferation was typically evaluate d on the third day via MTT assay. MTT assays were also performed on the first, sec ond, fifth and eighth day to better investigate possible cytotoxic effects of the substrate materials used. In each MTT assay the media was reduced from the original value of 2 ml to 0.5 ml and 75 l of MTT reagent were added to each well. After the yellow tetrazolium MTT reagent was completely converted into intracellular purple formazan by the me tabolically active cells (required time ~ 2 hours), all the media was removed from each well and the semiconducting substrates were transferred to new wells. In order to solubilize the formazan, 0.3 ml of dimethyl sulfoxide (DMSO) were added to each well and the absorbance of the solutions was determined spectrophotometrica lly with a plate reader (Bio Kinetics Microplate Reader EL 340, Bio-Tek Instruments) operating at a wavelength of 595 nm. Readings were corrected for the formazan formation due to the semiconducting surfaces alone. It should be pointed out that all the num eric values reported in the text above are for the wells containing semiconducting samples; values fo r the control wells were appropriately


62 scaled. The MTT assays were repeated three tim es for all the cell line s and, as explained above, performed in triplicate for the contro ls and for each different substrate. The obtained results are reported in the next s ection as sampling distri bution of the mean ( x) standard error of the mean ( m) and normalized with respect to the control readings. The adherent cell morphology was studied using fluorescence microscopy. Semiconducting samples and control wells were plated as described above and the cell morphology was inspected after 4, 24, 48, 72, 96, 120 and 216 hours (the latter only for BJ and HaCaT cells). Two hours before the optical inspection was performed using a Leica DM IL inverted microscope, the cel ls were fluorescently labeled with 2 L of CMFDA (5-chloromethylfluorescein diacetate) cell tracker dye. 3.1.3. SiC superior biocompatibility: MTT and fluorescent microscopy results MTT assays and fluorescent microscopy were used to evaluate the biocompatibility of the different SiC polytypes. Cell proliferation and adhesion quality of cells cultured on SiC surfaces were compared to the ones of cells cultured on control wells (which represent an ideal surface for cell grow th). Adopting the same procedure, SiC biocompatibility was compared to that of Si, which is at present the leading crystalline semiconductor for biotechnological applications The histogram in Fig. 3.3 reports the results of MTT assays performed on the thir d day for three different mammalian cell lines and clearly shows that SiC is, in all its pha ses, a high-quality surface for cell culture with significantly better performance than Si. No statistically significant differences were found among the cell proliferation on differe nt SiC polytypes. An extremely satisfying proliferation was observed for BJ cells on SiC, in average the same obtained for the


63 culture-well readings. In this case, a slight bi oactivity of the SiC substrate could even be hypothesized. The smaller difference in B16 cell proliferation observed between Si and SiC substrates can be easily justified by the cell line nature. B16 melanoma cells are in fact extremely aggressive cance r cells capable of indifferently adhering to substrates of diverse biocompatibility: therefore a reduced selectivity was expected. Figure 3.3. Cell proliferation of B16, BJ and HaCaT cells expressed as x mmeasured via MTT assays at the third day.Note that cell proliferation is greater on SiC than on Si surfaces for all the cell lines studied. As mentioned in the previous section, MTT assays were also perf ormed at the first, second, fifth and eight h day of culture. Cell seeding values for this particular set of assays were lower than the ones reported in § 3. 1.2 to avoid cell conflu ence before the eighth day of culture. The results obt ained for MTT assays performed on the first and eighth day of culture of HaCaT cells are reported in Fig. 3.4.


64 Figure 3.4. HaCaT cell proliferation on Si a nd SiC substrates measured via MTT at the first and eighth day of culture expressed as x m. Note that cell proliferation is greate r on SiC than on Si surfaces af ter eight days of culture. The MTT readings were approximately the same for all the samples studied (Si, SiC, and control wells) afte r 24 hours from seeding. This result suggests that no preferential initial adhesion is observed on SiC samples compared to Si and that, apparently, all the s eeded cells adhered to the semic onducting substrates. On the other hand, the cell proliferation at the eight day of culture was significantly lower on Si substrates than on SiC: calcula tions showed the reduction to be, with respect to SiC, as high as 56%. It should be noted that at the eighth day the cells reached confluence on the SiC substrates and control wells. The optical inspection results obtaine d using fluorescence microscopy supported the MTT quantitative results while providing additional useful information regarding cell morphology and adhesion quality. Fluorescent mi croscopy revealed that all the seeded cells were well distributed on SiC surfaces after 4 hours of incubation, elongated and flattened against the surface after 24 hours, and confluent (for the seeding density reported in § 3.1.2) after 4 days (B16 and Ha CaT), and 5 days (BJ). On the other hand, a lower density and generally inferior morphol ogy were observed for cells cultured on Si


65 substrates. Fluorescence microscopy images of adherent cells on Si a nd SiC substrates at the third day of culture are reported in Figures 3.5 and 3.6 at lower and higher magnification, respectively. Figure 3.5. Morphology of B16 ((a), (d)), BJ ((b), (e)) and HaCaT ((c), (f)) cells at the third day of culture on SiC and Si subs trates, respectively. Images by flourescence microscopy [88]. As is evident from images (a), (b), and (c) in Figures 3.5 and 3.6, cells on SiC substrates were well-distribut ed and flattened (the maximiza tion of the cont act area with the substrate is always an indicator of good biocompatibility) at the third day of culture. Moreover, their morphology was found to be identic al to that of the cells in the control wells: stellate for B16, elongated for BJ fibr oblasts and cuboidal for HaCaT cells. Even when imaged on the fifth or eighth day of cu lture (in experiments with lower cell seeding density), all the cells cultured on SiC substrates displayed an excellent morphology. This


66 finding, together with the fact that the cells could be cultured to c onfluence on all the SiC samples, confirms the lack of cytotoxicity of SiC substrates at least within the first eight days of culture. Figure 3.6. Higher magnification images of the morphology of B16 ((a), (d)), BJ ((b), (e)) and HaCaT ((c), (f)) cells at the thir d day of culture on SiC and Si substrates, respectively. Images by flourescence microscopy. On the other hand, the morphology of the cells cultured on Si substrates appeared to be inferior: typically a reduction in the cel l dimension was observed for all three cell lines. The reduced contact area of the adhere nt cells (low adhesion force) confirms the lower biocompatibility of the Si substrates. Specifically, B16 cells appeared rounded with a tendency to form clusters when cultured on Si (image (d) in Figures 3.5, 6). HaCaT and BJ cells were, at least during the first three days of culture, morphologically similar to the control cells but smaller and of a lower density (images (f) a nd (e) in Figures 3.5, 6). In


67 particular, the morphology of BJ cells culture d on Si degenerated at the fifth day of culture (Fig. 3.7(b)), while high-quality cel l morphology was observed on all of the SiC surfaces (Fig. 3.7(a)). The observed morphol ogical degeneration on Si surfaces was most probably caused by a not-null level of cy totoxicity of the Si substrates. Figure 3.7. Cell morphology of BJ fibroblasts at the fifth day of cu lture on a (a) SiC an d (b) Si substrates as measured with fl ourescence microscopy. Clearly the optimal morphology was observed on SiC s ubstrates due to the uniform distribution and elongate shape of the adherent cells. Another semiconducting material whose cyto toxic effect is well known is gallium arsenide (GaAs) [89, 90]. Since its deleteriou s effects on cells are well-known, we used it as a negative control to demonstrate the va lidity of our measurements. MTT assays were performed on GaAs and 6H-SiC samples af ter two days of B16 cell culture and the obtained results, reported in Fig. 3.8, clearly show the toxicity of this substrate material while again confirming the supe rior performances of SiC.


68 Figure 3.8. B16 cell prolifera tion measured via MTT at the second day of culture on 6HSiC and GaAs substrates and expressed as x m.Note that cell proliferation is greate r on SiC than on GaAs which is known to be cy totoxic (GaAs used as a negative control i n this experiment). The fact that even the aggressive B16 melanoma cells did not proliferate and stay viable when cultured on GaAs substrates conf irmed the high toxicity of this material and the validity of the procedure used to as sess the biocompatibility of SiC and Si. Fluorescent images of cells cultured on Ga As substrates showed a morphological degeneration of the cells (see Fig. 3.9(b)). Like in the case of Si, cytotoxicity provokes cytostructural changes and even tually leads to cell death. Figure 3.9. Morphology of hea lthy B16 cells cultured on 6H-S iC (a) vs. cytostructural degeneration of B16 cells cultured on GaAs as measured with fluorescence microscopy.


69 In conclusion, the MTT and fluorescence mi croscopy results show the superiority of SiC as a bio-surface for cell-culture while confirming a not-null level of cytotoxicity for Si. Also, the fact that different SiC pol ytypes, which display different band gaps and different doping levels (i.e., free charge density), performed in a similar way as substrates for cell culture suggests that the influence of the electronic propertie s of the material on cell adhesion and proliferation is not relevant in this case (at least for the SiC material system). However in Chapter 4 we will disc uss more carefully the electronic interface between cells and semiconductor surfaces and fi nal judgment of the influence of material electronic properties on cell viabilit y should be reserved till then. 3.1.4. SiC vs. Si: evaluation of cell pr otrusions via AFM and optical microscopy Additional studies were performed on cells cultured on SiC and Si surfaces using AFM and optical microscopy to quantitatively describe the presence of any cell extensions such as filopodia and lame llipodia. Lamellipodia are broad flat cell protrusions which surge forward and adhere to surfaces, allowing cel ls to gain traction and move on surfaces. Filopodia are rod-lik e cell surface projecti ons extending several micrometers ahead of the cells where they e xplore the extracellular surface. The presence and extension of filopodia and lamellipodia indi cate the quality of the cell response to a substrate material. Typically, numerous and extended filopodia and lamellipodia are indicative of good substrate biocompatibili ty. Therefore their measurement and quantification is an additional way to assess a material’s degree of biocompatibility. For these experiments B16 melanoma cells were seeded at a density of 40x103 cells/cm2 on 6H-SiC(0001) and Si(001) samples as described in § 3.1.2. After 24 hours the cells were fixed in 4% PFA (Paraformald ehyde) for 20 minutes at room temperature


70 (RT) and successively in a series of etha nol solutions with increasing concentration (17%, 35%, 70%, and 95%) for 2 minutes each. After the fi xation, platinum-gold was sputtered on the cell / surfaces in a gold deposition chamber as is typically done to prepare cells for scanning el ectron microscopy (SEM) analysis. The cells were then imaged using the XE-100 AFM system and its resident optical microscope at the PSIA facility in Santa Clara (CA). The optical inspection revealed, confirming wh at was described in § 3.1.3, a greater cell spreading and prolif eration on the SiC substrates. Cells on SiC appeared to be stellate in shape and interconnected through a bridging network (Fig. 3.10(a)). Cell bridges were found to be typically 5-10 m wide and 500 nm high (Fig. 3.10(b)). Typically, cells appeared to have reduced dimensions on Si substrates (as evident by comparison between Figures 3.10(c) and (e)). In general, extended lamellipodia, which indicate the formation of stable adhesion contacts, were observed on SiC surfaces (Fig. 3.10(d)). Instead, cells on Si had typically smaller or no lamellipodia present on their surfac es (Fig. 3.10(f)). Filipodia were present on cells cultured on both substrates and were typically 2-4 m wide and 5-15 nm high as reported by the line pr ofile in Fig. 3.10(f). Note that the two holes present in the cell in Fig. 3.10(f) are most likely nuclei in the mitotic phase which exploded during the gold sput tering process, which is done under vacuum. The missing nuclei allow one to see some beaded features 1 to 2 m in diameter which are probably, because of their central location, fibrillar adhesions of the cells to the substrate.


71 Figure 3.10. Optical (left) and AFM (right) mi crographs of (a-d) B 16 cells on SiC and (e, f) Si surfaces. In particular : (a) the observed bridging networ k of cells with (b) ‘bridge’ dimensions; cells presenting (c) extend ed lamellopodia on SiC and (d) single lamellopodia with relative dimensions; (e) ce lls presenting filopodia on Si and (f) a single cell with relative filopodia dimensions. AFM micrographs (b), (d) and (f) are 3030 m, 4040 m and 4545 m, respectively. 3.2. Influence of surface properties on cell adhesion and proliferation The previous section reports an evident gr eater biocompatibility of SiC with respect to Si. In § 3.2.1 we search for the reasons at the basis of the observed greater biocompatibility by investigating the chemis try and the wettability of SiC and Si


72 substrates. Sections 3.2.2 and 3.2.3 focus on the effect that the topography and chemistry of SiC surfaces have on cell proliferation, respectively. 3.2.1. Surface chemistry and wettability as possible explanations of SiC greater biocompatibility It has been already pointed out that the SiC and Si samples used in § 3.1 present similar topographical characteri stics: therefore substrate topo graphy can be excluded as a cause of the reported SiC superior biocompatib ility with respect to Si. We identify three main remaining reasons that could be the basis of the enhanced cell proliferation observed on SiC substrates: 1) influence of the different surface chemistry on cell adhesion / proliferation; 2) differences betw een Si and SiC surface wettability. 3) reduced electronic interaction (e.g., charge exchange ) between cell adhesi on proteins and SiC surfaces (SiC has a larger energy bandgap, Eg, than Si and therefor e a lower density of unoccupied states in the energy rang e of the transfer level, i.e. Eg > 1.8 eV [30]). Since (3) is a rather complicated issue which ma y find an answer only after the electronic interaction between semiconducti ng substrates and biological ce lls will be exhaustively understood (and therefore goes beyond the purpose of this work), we investigate (1) and (2) here with preliminary studies pert aining to (3) presen ted in Chapter 4. For this purpose, XPS and contact angle st udies were performed on the same Si and SiC surfaces used in the MTT assay described in § 3.1.3 after they were cleaned as described in § 3.1.1. Representative XPS surv eys of as-introduced Si and SiC samples are shown in Fig. 3.11.


73 Figure 3.11. XPS surveys of two of the Si and SiC samples used in § 3.1.3.No comparison has to be made between the pe ak magnitudes of the two samples; only elemental percents within an individual sample are significant. Note the expected greate r C-concentration on the SiC surface. From the chemical analysis of the near surface region it app ears that both Si and SiC present the same elements at the surface, namely Si, C, and O. However, the (graphitic) carbon present on Si is in lowe r concentration (typical ly < 25%) than the carbon on SiC (typically > 50%). The higher Cconcentration on SiC surfaces is a likely explanation for the greater biocom patibility of SiC. It is in fact well-known that surfaces with a normal electrochemical potential cl ose to the one of the cells are more biocompatible [11]. Since the el ectrochemical potentia l of carbon is comparable to that of living tissue, the higher carbon c oncentration of the SiC surfaces most likely justifies SiC greater biocompatibility. Wettability of Si and SiC samples was ev aluated using the sessile drop method at the IFN-CNR of Trento, IT [91]. A 1 L dropl et of DI water was deposited on the surface of each analyzed sample by using a biologi cal pipette. The droplet contact angle ( ) was imaged using a contact angle goniometer and estimated by measuring the angles between


74 the baseline of the drop and the tangent at the drop boundary using the ImageJ 1.37v software. At least two droplet s were deposited on different areas of each sample. The analysis of each droplet produced two contact angle values (e.g., left and right). Therefore the contact angle value was calcula ted for each sample as the mean of four values. Contact angle images for the SiC and Si samples analyzed are reported in Fig. 3.12. The contact angle values reported in Table 3.2 are expressed as the sampling distribution of the mean ( ) the standard error of the mean ( m) from a three samples distribution for each substrate material / polytype. Figure 3.12. Measured water contact angles on Si, 3C-, 4H-, and 6H-SiC surfaces.Note the higher hydrophobicity of Si. In this e xperiment surface roughne ss does not influence the contact angle values since all the samples display similar values of Rq. Table 3.2. Wettability of SiC and Si surf aces as measured via sessile drop method. Substrate Material Si(001) 3C-SiC(001) 4H -SiC(0001) 6H-SiC(0001) () () 88.28 1.92 28.32 2.27 20.27 3.00 25.2 2.13 From the images and the values reported we can infer that Si surfaces are considerably more hydrophobic than SiC surfaces. No relevant differences were found in the wettability level of the different SiC pol ytypes studied. As demonstrated by the tight standard deviation values, the surface wettability properties of the different samples used in the MTT assays in § 3.1 were similar for each substrate material / polytype. The high level of hydrophilicity exhibited by the SiC su rfaces may also be a likely explanation for


75 its better performance as a substrate for cell culture. In fact, even though cell adhesion proteins adhere better to hydrophobic surfaces, mammalian cells are known to preferentially adhere to hydrophilic surfaces [92]. In conclusion, the lower biocompatibility of Si could be due to a competitive mechanism where both the electrochemical potential (e.g., absence of carbon from the surface) and the surface energy (e.g., higher hydrophobicity) play an important role. It has also to be pointed out that Si exhibits much poorer tribological properties than SiC, which means a higher density of debris forma tion. Cell phagocytosis of the toxic debris may likely lead to the cytostructural degenera cy reported for cells cu ltured on Si for more than 4 days (see Fig. 3.7). 3.2.2. Influence of SiC surface topogr aphy on cell adhesion and proliferation Several studies report that surface roughness influences cell adhesion and migration. Typically, higher values of root mean square roughness yield higher osteoblast proliferation [93, 94], while flatter surfaces have been shown to enhance fibroblast proliferation and adhesion quality [95]. For this reason we investigated the influence that varying surface roughness of SiC may have on cell morphology and proliferation. For this purpose we selected SiC samples with different surface charac teristics: three asgrown (Fig. 3.13(a)) and three over-etched (F ig. 3.13(b)) 3C-SiC samples whose surface features were exhaustively described in § 2.2.1. As a point of reference Rq was roughly 1.4 nm for a 25 m2 area of the as-grown samples and 35 nm for the same area of the over-etched samples. The substrates were clea ned as described in § 3.1.1, which excludes any possibility of H-termination due to the etching process on the over-etched surfaces. Both B16 melanoma cells and BJ fibroblasts were cultured on those substrates. MTT and


76 optical inspection were performed at the third day as described in § 3.1.2. The results of both analyses showed no difference in cell proliferation and adhe sion quality on the two different sets of samples. Figure 3.13. AFM micrographs showing the mo rphology of the 3C-SiC samples used to evaluate the effect of surface roughness on ce ll adhesion and proliferation. As-grown (a) and over-etched (b) 3C-SiC with Rqvalues of 1.4 nm and 35 nm, respectively. AFM data taken in tapping mode. This result may be explained by the fact that the roughness variation between the two types of 3C-SiC surfaces studied was one to two orders of magnitude smaller than the cell adhesion plaques which are responsible for cell adhesion (34 nm of surface roughness variation against the 500-3200 nm of th e adhesion complexes [96]). A study of Zhu et al. reports that patterned s ubstrates need to have features ranging between 100 and 3000 nm in order to influence the plaque dynami c and therefore the ce ll adhesion pattern [96]. To verify this finding, BJ fibroblast cells were cultured on a 3C-SiC sample whose edge, after it was accidentally slivered, disp layed macroterraces which were 2000 to 4000 nm wide. Optical inspection was performed af ter 24 hours and showed that the cells had spread in a patterned fashion within the slivered edge. The majority of the fibroblasts were well-spread and elongated along the macrosteps as shown in Fig. 3.14.


77 Figure 3.14. Patterned fibroblasts adhesion on the slivered edge of a 3C-SiC sample which displays macroterraces 2000 to 4000 nm wide. Image by fluorescence microscopy. Summarizing, surface micropatterning s eems to have a bigger effect on the morphology of fibroblast adhesion than surf ace nanopatterning. Variations of roughness in the nanometric scale seem to have little or no effect on the morphology and proliferation of B16 melanoma cells and BJ fi broblasts. This result is in agreement with what was reported by Richards [97]. Nonetheless, atomically flat surfaces like the ones produced via H-etching and described in Chapter 2, may be preferable to the commercially polished ones not only for biomolecular surface science studies but possibly also for implantable biosensors, since their smoothness reduces the chance of bacterial adhesion and therefore infection (e.g ., common infectious bacteria diameter ~ 2m). 3.2.3. Influence of SiC surface chemis try on cell adhesion and proliferation Multiple studies have shown that the su rface chemistry of bi omaterials strongly influences cell adhesion and proliferation a nd may regulate the cellular behavior at the molecular level [98, 1, 18]. A previous work conducted on polyacrylamide surfaces has


78 shown a preferential adhesion of mamma lian cells on hydroxylic group (-OH) and carboxylic acid (-COOH) term inations [99]. Curtis et al. report that hy droxylic groups considerably increase cell attachment on a la rge number of polymeric surfaces such as polypropylene and polystyrene [100] We initially thought that H2-etching of SiC samples may have terminated the surfaces in such a fashion. Since the AE S studies reported in Chapter 2 were inadequate to detect hydroge n on the etched surfaces, we performed total attenuated reflectance Fourier transf orm infrared spectroscopy (ATR-FTIR) measurements, which we describe in Chapter 4, to probe whether –OH groups are present on H-etched 3C-SiC. Indeed, the ATR-FTIR results indicated the presence of C-H bonding at the surface while –OH groups wh ere detected on regularly HF dipped surfaces. To probe whether the different surface terminations had an influence on cell proliferation we cultured HaCaT cells on H2-etched and HF treated 3C-SiC samples and observed their behavior. Cell proliferati on was evaluated through MTT assays while culture wells were used as controls. It has to be pointed out that bot h the H-etched and the HF treated samples were 810 mm diced from the same 2 inch wafer and therefore presented comparable surface characteristics. The etching was performed for 5 minutes at 1200 C, at atmospheric pressure and with a hydrogen flow equal to 10 SLM in the CVD reactor described in § 2.2. H-etching di d not produce any change of the surface morphology of the 3C-SiC samples while hydr ogen was detected on the surface by ATRFTIR analysis. The etched samples were not cleaned to preserve the H-termination while the un-etched samples were cleaned as described in § 3.1.1. The MTT assay was performed at the third day and in triplicat e for the etched and HF treated 3C-SiC


79 substrates and the controls. Th e obtained results are reported in Fig. 3.15 and expressed as x m. Figure 3.15. Comparison in HaCaT cell prolif eration between etched and un-etched 3CSiC substrates at the third day of culture expressed as x m measured via MTT assay. Note no difference in proliferation was obser ved despite the detection of hydrogen on the etched surface via ATR-FTIR. As is evident from the histogram above, ce lls proliferated at the same rate on Hetched and un-etched 3C-SiC surfaces. Fluo rescence microscopy confirmed this result which may be explained considering that the elapsed time between cell seeding and adhesion is roughly four hours a nd that during this time numerous chemical changes take place at the semiconductor surface. More proba bly the surface oxide layer, caused by the immersion in media, is formed before the cel l attachment is finalized. At that point the un-etched and etched surfaces will present mo st likely similar chemical characteristics (i.e., any hydrogen effect on the surface is eliminated by the oxide). We already discussed that a higher Cconcentration on the SiC surfaces is a possible cause of the SiC greater biocompatibility and that, in general, a C-rich surface is believed to enhance cell adhesion. For this reas on we decided to study the proliferation of HaCaT cells cultured on the C-face of SiC sa mples. In this experiment three 4H-


80 SiC(0001) and\ three 4H-SiC(000 1) surfaces, nominally Siand C-face, respectively, where used as substrates for cell culture afte r being cleaned as described in § 3.1.1. No differences in surface morphology were observed among the different samples. The MTT assay, performed at the third day of cell culture, showed a surprisingly reduced cell proliferation on the C-face samples (Fig. 3.16). Figure 3.16. HaCaT cell proliferation at the third day on Siand C-face 4H-SiC expresse d as x m measured via MTT assay. Higher pro liferation is observed on the Si-face. However, light was shone on this puzz ling result by XPS anal ysis performed on 4H-SiC(0001) and (000 1) substrates cut from the same wafer and cleaned at the same time as the samples used for the MTT assay. Surface elemental ratio values, as calculated from the surveys, are reported in the first thr ee rows of Table 3.3. The percents of C and Si elemental concentration on the surfaces, reported on the last two rows, are calculated from the high resolution scans.


81 Table 3.3. Elemental concentrations a nd ratios for 4H-SiC(0001) and 4H-SiC(000 1). 4H-SiC(0001) Si face 4H-SiC(000 1) C face C/Si 1.62 1.71 O/Si 0.34 0.35 F/O 0.33 1.30 C % HR 60.05 63.14 Si % HR 39.95 36.86 While the values in Table 3.3 confirm th e higher C-concentration of the C-face surfaces, the most striking result is the hi gh concentration of fluorine (F) found on 4HSiC(000 1). It has to be noted that such a hi gh concentration of F was never found on all the other SiC (always the Si-f ace was studied) surfaces examin ed in the past after being cleaned as in § 3.1.1. Low F concentrations were found rarely and were mostly caused by a shortened rinsing time after the HF dip. Si nce both fluorine and fluor ide ions are toxic, a high F-concentration on the 4H-SiC(000 1) is a likely explanation for the reduced cell proliferation observed on those surfaces. However, at this time it is unclear why the C-F bond should be more resistant to a DI water rinse than the Si-F bond. Fluorine is known as the most highly electronegative element on ea rth (e.g., 3.95 in th e Pauling scale), and will therefore create stronger bonds with elements exhibiting a lower electronegativity. Si electronegativity is lower than that of C (1. 90 versus 2.55 in the Pauling scale): therefore the silicon-fluorine bond, which pr esents an energy of 582 kJ mol-1, is extremely stable and evidently stronger than the carbon-fl uorine bond, whose energy is 452 kJ mol-1. Therefore, a higher fluorine residue would not be expected on C-face samples after DI water rinse. However, Jacobsohn et al. [ 101] report that the incorporation of CFn groups


82 on the surface of SiC make it much more hydrophobic while the pres ence of Si-F bonds doesn’t change the material wettability. Therefore, the way the fluorine atoms are incorporated on the SiC surface may play an important role: the increased hydrophobicity on the C-face surfaces could likely reduce the ef fectiveness of the DI rinse step. Also, a lower fluorine concentration on Si faces imme diately after HF exposure and prior to the DI rinse step could be another possible expl anation for the reduced presence of fluorine in Si rather than in C faces. In fact, Si su rfaces have been found to present mostly silicon hydride species and very little oxide or fluorid e after HF treatment [102]: because of the high electronegativity of fluorine, Si-F bonds are prone to polari ze Si atoms which may become much more susceptible to nucleophilic attack than Si-H bonds. In any case, the DI rinse step duration should be increased for 4H-SiC(000 1). 3.3. Cleaning of SiC surfaces for bio-applications: RCA vs. Piranha The previous section has drawn attention to the importance of choosing appropriate cleaning procedures when preparing surfaces fo r biological evaluation. In this section we compare the effectiveness of two standard cl eaning procedures used to remove organic residue from SiC surfaces prior to cell growth : RCA (Radio Corporation of America) and Piranha cleaning. The efficacy of these two procedures was evaluated through optical and chemical analysis, and Piranha was shown to be the one capable of completely removing the organic contaminants (residue after cell expo sure) and, therefore, best able to produce chemically reproducible surfaces. However, th e ineffectiveness of RCA cleaning, which leaves bio-residue on the surface, may be us eful for identifying the chemistry and shape


83 of the cell focal adhesions (§ 3.3.2). Thus in § 3.3.3 we report on the effects of repeated Piranha cleanings on surface chemistry and wett ability, which affect cell proliferation. 3.3.1. Effect of RCA and Piranha clean s on semiconductor surface morphology and chemistry A total of six 4H-SiC(0001) an d four 3C-SiC(001) samples, all of them presenting atomically flat surfaces and constantly used as substrates for cell growth, were used to evaluate RCA and Piranha effectiveness in organic residue removal. The methodology adopted was the following: first, HaCaT cells were grown on the surfaces of those samples; second, half of the samples were cl eaned with Piranha and half with RCA; last, their surface morphology and chemistry was evaluated. The Piranha procedure adopted was the one described in § 3.1.1. The step s of the RCA cleaning procedure were as follows: immersion in a 1:1:6 mixture of NH4OH:H2O2:H2O heated to 65 C for 10 minutes; rinse with DI water; dip in a 50:1 H2O:HF for 30 seconds; rinse in DI water; immersion in a 1:1:6 mixture of HCl:H2O2:H2O heated to 65 C for 10 minutes; rinse in DI water. The oxide generated by the oxidi zing steps of RCA was removed and the samples were sterilized by using the same procedure adopted after Piranha clean and described in § 3.1.1. Which means, di p in hydrofluoric acid solution (50:1 H2O:HF) for 2 minutes; thoroughly rinse in DI water; ethanol dip followed by DI water rinse; nitrogen dry. The surface morphology of the samples was inspected using the optical camera of the XE-100S PSIA AFM system.


84 Figure 3.17. Optical images showing the su rface morphology of 4H -SiC(0001) after cell growth and subsequent cleaning with (a) Piranha and (b) RCA cleans. Note the highe r level of residue after RCA cl eaning indicating that Piranh a is superior in producing chemically repeatable surfaces. Fig. 3.17(a) and Fig. 3.17(b) show the surface morphology of 4H-SiC(0001) cleaned with Piranha and with RCA, respect ively. While the surface of the Piranhacleaned sample appears completely spotless the surface of the RCA-cleaned sample is heavily contaminated by organic residue. It su rprisingly appears that even dead cells are present on the surface after RCA cleaning (see ce nter of Fig 3.14(b)). Similar results were obtained for the 3C-SiC surfaces. The surface chemistry of the samples was analyzed via XPS. XPS surveys and high resolution scans of the RCA cleaned samples displayed spectral lines whose kinetic energy was clearly associated to nitrogen (N) and sodium (Na) (Fig. 3.18), while elements other than C, Si and O were not de tected in samples cleaned with Piranha. Since N is the key atom in the amines of the adhesion proteins, its presence indicates the existence of proteins on the surface.

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85 Figure 3.18. Presence of nitrogen (N) and s odium (Na) in the XPS survey of 3CSiC(0001) after cell growth and subsequent RCA clean.These elements are indicators o f residual biomatter on the surface and N, in pa rticular, is a key component of adhesion proteins. The reported results lead to the conclu sion that Piranha clean is much more effective in removing organic residue than RCA clean. For this re ason, Piranha solution was constantly used in the sample cleaning procedure prior to cell deposition (§ 3.1.2). 3.3.2. RCA clean as a promising surface pre-treatment for the study of cell-SiC adhesion sites When an epithelial cell adhe res to a surface three typical types of adhesions may take place: focal complexes, focal adhe sions and fibrillar adhesions, which are characterized by different molecula r constituents and dimensions. In-vivo, these adhesions exist between the cells and the extra-cellular-matrix (ECM), which is a fibrous mesh of proteins that serves as both structur al scaffold and substrate for the display of signaling ligands. At present, many studies are focusing on rebuilding ECM surrogates in synthetic materials to engin eer cell-surface inte ractions [96, 103, 104]. However, when cells come into contact with a substrate th at naturally does not present ECM proteins,

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86 they may or may not produce ECM depending on the substrate level of biocompatibility. Formation and secretion of ECM proteins, incl uding fibronectin and type I collagen, are required for cell spreading and anchorage, so biocompatible material will enhance ECM production. We used the presence of bio-residuals on the surface after RCA cleaning to study the nature of the HaCaT a dhesion sites on SiC surfaces The identification, by topographical exam, of the typology of a dhesions could help identify the adhesion proteins that have an active role in SiC-HaCaT cells interaction. Typical optical images of RCA-cleaned SiC su rfaces were observed to be similar to the one of Fig. 3.19. Dead cells are present in the centre of the im age, while dot-shaped particles, which are possible adhesion sites, ar e visible in the right hand side (RHS) of the figure. The height of the cell residue was m easured to be 200-500 nm (Fig. 3.19(b)) by AFM. Also, elongated fibrillar elements, whose nature is explained further on in this section, were observed frequently (see upper LHS of Fig. 3.19(b)). Figure 3.19. Analysis of RCA cleaned su rface. (a) optical micrograph and (b) AFM micrograph of a cell observed in (a) with cell dimensi on shown via the relative line profile.

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87 A better analysis of the dot-shaped featur es can be performed from Fig. 3.20(a) which probably represents (considering the s cale bar and that the size of an extended HaCaT cell is ~ 20 m) a site for multiple cell adhesion. Figure 3.20. Analysis of different adhesi on sites found on SiC surfaces after RCA clean. (a) optical micrograph and (b) AFM micrograph of a particular feature in (a) depicting what are believed to be podosomes on the surface. We identified, by analyzing their dimens ions, three different typologies of cell adhesions that may still be present on the Si C surface after RCA cleaning (see labeling in Fig. 3.20(a)). Fibrillar adhesions are beaded features of variable dimensions ranging from 1 to 10 m. Smaller dot-like features are e ither focal complexes (~ 1 m) or podosomes (<1 m, Fig. 3.20(b)) [105, 106]. In many samples a fibrillar network like the one in Fig. 3.21 was observed. This suggests massive ECM production on SiC substrat es, which also confir ms the material’s biocompatibility. In the ECM net were often observed circular shaped holes, which probably indicate the original pr esence of cells on those sites (s ee circle in Fig. 3.21(a)). In fact, if the ECM net-cell is a stronger bond than ECM netSiC substrate, when the cell is detached by the chemical cleaning, the fibrillar net will also be removed. AFM

PAGE 103

88 micrographs and related line profiles, like the ones in Fig. 3.21(b), reported the fibrillar elements to be up to 300 nm high and 0.5-1.5 m wide. Figure 3.21. (a) Optical image and (b) AF M micrograph, with related line profile, showing a fibrillar network on the SiC surf ace after RCA cleaning. The circle in (a) indicates an area of missing fibers and theref ore a possible cell site before RCA cleaning. Cellular residue are visible in the upper RHS and lower LHS of the optical image. 3.3.3. Effect of repeated Piranha clean ings on chemistry, wettability and cell proliferation While performing the biocompatibility studi es whose results are reported in § 3.1 and 3.2, we noticed that the cell proliferation was signifi cantly lower on samples cleaned with Piranha several times (e.g., more than ten) than for samples cleaned only a few times. As a consequence, we decided to investigate if this difference in cell proliferation was caused by changes in surface properties in duced by the Piranha clean chemistry. For this purpose, XPS studies and contact angle an alysis were performed on a series of 3CSiC(001) and 4H-SiC(0001) samples cleaned in Piranha from a minimum of zero to a maximum of twelve times. Fig. 3.22 reports the HaCaT cell prolif eration on samples never treated with Piranha a nd treated ten times, respectively. The results, calculated from a series of three MTT assays, showed that the cell proliferation for samples never treated with Piranha was up to 23% highe r than for the over-treated samples.

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89 Figure 3.22. HaCaT cell proliferation at th e third day on 3Cand 4H-SiC samples neve r cleaned with Piranha and cleaned w ith Piranha 10 times expressed as x m measure d via MTT assay. [88] Contact angle studies performed on the diffe rently treated 3Cand 4H-SiC surfaces gave the results listed in Table 3.4. Table 3.4. Piranha effect on surface we ttability assessed by sessile drop method. # piranha treatments3C-SiC(001) 4H-SiC(0001) 0 52.53 1.83 53.78 3.38 3 35.2 1.7 () () >10 22.8 2.1 20.65 3.35 It appears that Piranha cleaning tends to increase the surface hydrophilicity. By comparing the results in Table 3.4 and Fig. 3.22, we can state that in this case, differently than in § 3.2.1, a higher hydrophilicity does not correspond to a greate r cell proliferation. This apparently contradictory result (e.g., mammalian cells should preferentially adhere and proliferate on more hydrophilic surfaces) was explained by the results of XPS analysis. In Table 3.5 we re port the data extracted from XPS survey and high resolution scans of 4H-SiC sample never treate d with Piranha and treated 10 times.

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90 Table 3.5. Si and C elemen tal concentration* and C/Si ratio** for 4H-SiC samples treated with Piranha zero and ten times. as calculated from the XPS high resolution scans ** as calculated from the XPS survey No Piranha Piranha 10 times C/Si ** 1.62 1.56 C % 60.05 56.28 Si % 39.95 43.72 It must be noted that th e 4H-SiC surfaces examined with XPS were of vicinal samples cut from the same two inch wafer and therefore presented similar chemical characteristics before the Piranha treatment. It is evident in Table 3.5 that repeated Piranha cleanings reduce the C-concentration on the near surface region in a sensitive way. This result suggests that, even though SiC is well-known to be chemically inert to the most widely used acids and liquid etchants, it is likely that repeat ed piranha cleans of SiC surfaces may decrease the carbon concentr ation in the near surface region. Piranha is, in fact, a strong oxidizer and highly effective in removing or ganic matter. Its extremely reactive atomic oxygen species, that form during H2O2 dehydration, allows the Piranha solution to dissolve elemental carbon, which is notoriously resistant to room temperature aqueous reactions. A decreased C-concentration in the Piranha over-treated surfaces is a likely explanation for the increased hydrophilicity observed on these surfaces. Ab-initio molecular dynamic simulations reported in [107] have, in fact, shown that carbonterminated SiC surfaces have a hydrophobic char acter, while Si-terminated SiC surfaces are hydrophilic. Also, other sources report th at a surface covered with a substantial amount of carbon bonded to oxygen is hydrophobic [ 13]. Therefore it is likely to assume

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91 that a reduced C-concentration on the SiC su rface may lead to a consequently increased hydrophilicity. Also, since surfaces richer in carbon are more biocompatible, the reduced C-concentration observed for Piranha over-treate d samples is very likely the cause for the decreased cell proliferation reported in Fig. 3.22. Summarizing, the repeated use of Pi ranha cleaning changes the SiC surface chemistry and subsequently its wettability and biocompatibilty. In order to obtain consistent results in biocompatibility studies the Piranha cleans of a specific sample have to be limited to a maximum of five. This pa rticular precaution was used when performing all the MMT assays and optical in spections described in § 3.1.3. 3.4. Summary The presented studies report the significan t finding that SiC surfaces are a better substrate for mammalian cell culture than Si in terms of both cell adhesion and proliferation. In the past, the fact that cells could be direct ly cultured on Si crystalline substrates led to a widespread use of these materials for biosensing applications [108110]. Therefore, the resu lts reported in § 3.1 define SiC an even more promising substrate for future cell-semiconductor hybrid systems. The main factors that have been shown to define SiC biocompatibility are its hydrophilicity and surface chemistry (§ 3.2. 1). SiC surface morphology is shown to influence cell adhesion only when macropattern ed (§ 3.2.2), while SiC polytypism and doping concentration seem to have no infl uence on cell proliferation (§ 3.2.3). The importance of using an appropriate cl eaning procedure to obt ain repeatable and clean surfaces after each cell cult ure cycle is discussed in § 3.3.1 and in § 3.3.3. In § 3.3.2

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92 we show how the finding that RCA cleaning leaves bio-residuals on the surface can be used to analyze the properties and morphology of the adhesion sites. Many more analysis techniques may be us ed to analyze SiC biocompatibility. In particular, primary cell lines could be culture d on SiC surfaces in the future since their behavior would be a closer description of the in-vivo performance of the material. Also, since cell adhesion is known to be influenc ed by the electrostatic properties of the substrate in the specific media where cells ar e grown, zeta potential measurements of SiC particles in media could be attempted to defi ne its charge. The adhesion sites of cells on SiC surfaces could be analyzed by using specific dies wh ich would allow for a more accurate adhesion typology differentiation than the one presented in § 3.3.2. Once the adhesion sites are well characterized a mo re precise identifica tion of the adhesion proteins would be possible using XPS and FTIR. The iden tification of the organic chemical groups that bind to the SiC surface, together with the calculation of SiC zeta potential in media, could be used to better understand the electronic inter action between cell and SiC surfaces.

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93 Chapter 4. CPD Apparatus and Charac terization of Surfaces for Cell-Semiconductor Electronic Interaction Studies The surface characteristics and biocompatibility of silicon carbide (SiC) have been extensively studied in Chapters 2 and 3, re spectively. Thanks to the knowledge acquired through these studies we can now move fo rward and characterize cell-semiconductor electronic interactions via co ntact potential difference (CPD) measurements. To date, dark/light CPD measurements have been prove n to be a powerful technique which allows one to investigate the electronic band bandi ng naturally present at the surface of a semiconductor or which occurs as a result of chemical or ionic charging [40, 42, 45]. The measurements that we aim to perform and that will be extensively described in Chapter 5 intend to investigate the band bending induced in a semiconductor by the charge associated with biological ce lls residing on the semiconductor surface. However, there is a potential complication that th ese measurements intrinsically present: the necessity of a liquid layer (i.e., cultu ring cell media) around the cells which implies immersion of the semiconductor in an electrolyte. In order to perform successf ul CPD measurement of a semiconductor immersed in liquid, the CPD apparatus has to be carefully assembled and calibrated and the samples selected for the liquid measurements have to fulfill several requirements: (i) biocompatibility to allow cell adhesion during th e experiment with adherent cells; (ii) low surface state density to impede pinning of th e Fermi level in liqui d; (iii) complete

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94 electronic characterization in ai r via CPD. All the SiC surfaces fulfill requirement (i), as shown by Chapter 3, while we suggested in Ch apter 2 that a low surface state density (ii) can be obtained on damaged SiC surfaces by pr ocessing them with suitable H-etching processes. In this chapter we therefore concentrate on requirement (iii). After describing the CPD apparatus and its calibration, we present the experimental procedure designed to perform the CPD measur ements which are used to electronically characterize semiconducting surfaces (§ 4.1). Suitable samples for cell-semiconductor electronic interaction studies are carefully selected (§ 4.2) and electronically characterized via CPD measurements (§ 4.3-5). In particular their electronic ‘steady state’ is studied in § 4.3, while the effect of Hetching and selected chemical treatments is described in § 4.4 and § 4.5, respectively. 4.1. CPD apparatus, calibration and experime ntal procedure for CPD measurements in air The CPD apparatus was properly assemb led to reduce stray capacitances and electromagnetic, mechanical and noise interfer ence and calibrated to obtain the maximum possible accuracy of the instrument. § 4.1.1 de scribes the experiment al apparatus and its assembly while § 4.1.2 particularly focu ses on the apparatus ca libration and suggests several precautions that may be taken while performing CPD measurements. In § 4.1.3, we describe the classical procedure adopt ed when measuring the surface potential ( s) of a dry semiconductor.

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95 4.1.1. Experimental apparatus As explained in Chapter 1, the component s of a basic CPD system are a vibrating Monroe probe connected to an electrostatic voltmeter for the detection of CPD voltages and a grounded chuck for sample backside gr ounding. To facilitate the reader in the understanding of the next sections we reca ll from § 1.4 the following concepts: (i) a Monroe probe measures a voltage that is defined as the CPD voltage (Vcpd); (ii) in darklight measurements two different voltages ar e acquired by the CPD system, one in dark (Vcpd,dark) and one under deep illumination (Vcpd,light); (iii) the difference between these two values yields a semiconductor surface potential (e.g., s = Vcpd,dark – Vcpd,light). The Monroe probe used in our system is a 1017AEL-5 model from Monroe Electronics (ME) [111]. The 1017 AEL-5 probe presents a large diameter sensing aperture (i.e., 3.30 mm) which enables low noise measurements at a probe-to-surface spacing as large as 6 mm. Its output was conn ected to an SDI (Semiconductor Diagnostic, Inc.) PDM-40 electrostatic voltmeter which resolved the CPD voltage (Vcpd) sensed by the probe by applying the null-method suggest ed by Kelvin [44]. The same voltmeter provided the ground reference potential for the gold-coated vacuum chuck, which is the platform where the sample was positioned duri ng the measurements. In order to exclude external electromagnetic fiel d and possible noise sources th e probe and the chuck were situated inside a Faraday cage which was properly grounded. An image of the basic experimental apparatus inside the Fa raday cage is shown in Fig. 4.1.

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96 LED probe sample LED probe sample Figure 4.1. Experimental CPD apparatus hous ed inside a Faraday cage (probe, sample, chuck and LED shown). Note: probe is shown protruding out of the probe holder arm for display purposes only. The output of the PDM-40, which displays, when properly calibra ted, voltages with a maximum accuracy of 1 mV, was connected to a Keithley 2000 voltmeter which was interfaced to a computer for automated data collection through an IEEE GPIB interface card. A suitable program was coded in the measurement control software Labview for data collection and storage. The CPD voltage as output from the Labview program after each measurement was calculated from an aver age of 25 samples taken with a sampling period of 10 ms. A Tektronix TDS 1002B oscilloscope was us ed as a diagnostic aid for the PDM-40 calibration. In order to perform the dark/light measur ements necessary for the determination of the semiconductor surface potential two ultr aviolet (UV) light emitting diodes (LEDs) with a 370 nm primary emission wavelength were positioned within th e terminal part of the probe holder with an angle that al lowed proper illumination of the semiconductor surface for a fixed range of probe-sample distances (Fig. 4.1). To allow proper illumination of the sample the Monroe probe was used while retracted in the probe holder

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97 (e.g., a different arrangement than that show n in Fig. 4.1). UV light was chosen because the photon energy of the light needs to be gr eater than the SiC energy band gap (i.e. Euv > 3.23 eV) in order to generate enough el ectron hole pairs (EHP) for the semiconductor surface to approach flatband. 4.1.2. CPD system calibration and measurement precautions Since surface potential is highly dependent on the semiconductor material morphological and electronic pr operties, CPD measurements are sample variable and definitely non-trivial to make with a high enough level of precisi on to gain useful information. Therefore several precautions ha ve to be used when performing these types of measurements. In this section we descri be the calibration procedure and precautions which allowed us to perform good-quality CPD m easurements in air and that, in the end, made possible a successful characteriza tion of the cell-semiconductor electronic interaction. Measuring distance selection. Even though the AEL1017-5 Monroe probe model allows measuring distances as large as 6 mm the choice of operating at small probe-tosample distances was made based on the fact that larger distances make the measurement more susceptible to external noise and d ecrease the electrostatic voltmeter response speed. The PDM-40 voltmeter was periodically calibrated to ensure spacing-independent measurements for a probe-to-sample dist ance ranging from 0 to 2.5 mm. Even though this precaution was constantly taken we opted to work at a constant probe-to-sample distance to avoid the effect of stray capacitances, which are known to be distance dependent [42]. The optimal operating prob e-to-sample distance was selected to maximize both the sample illumination (e.g., ma ximization of the light intensity on the

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98 area below the measuring electrode) and the me asurement accuracy (e.g., minimization of the standard deviation of a series of r eadings from the same surface). While the maximization of the sample illumination was straightforward (we selected the range of distances for which the two di scs of light originating from the UV diodes would intersect on the sample surface), the determination of the measurement accuracy required us to take several Vcpd readings of the grounded chuck at different probe heights and to calculate the measurement standard deviat ion at each specific distance. Mean ( cpdV) and standard deviation ( ) for a three measurement distributi on at each analyzed distance are listed in Table 4.1. The probe-to-sample dist ances analyzed are the ones which allowed the maximization of the light coverage on the sample surface. Table 4.1 CPD voltage mean ( cpdV) and standard deviation ( ) measuredfor the grounded chuck in the dark and at di fferent probe-to-sample distances. Probe-to-sample distance (mm) 1 mm 1.6 mm 2.3 mm cpdV (mV) 0.42 0.24 1.08 1.34 0.28 1.96 As expected, the smallest standard deviat ion was obtained for the smallest probe-tosample distance which was therefore selected as the optimal measuring distance. Hence, all the measurements described in this chapter were performed with the probe 1 mm above the sample surface. Verification of the signal transient dete ction capability. Besides reading the CPD voltage of a semiconductor in equilibrium (e.g., Vcpd is constant since no photoexcitation, electric field or chemical charging of the surface is applied), the CPD system has to be able to detect quick variations in the CPD voltage wit hout significant delay.

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99 This is because when light shines on the se miconductor surface the voltage quickly shifts from Vcpd,dark to Vcpd,light. The capability of the apparatus to detect fast signal transients was checked by applying a square waveform to the metallic chuck using a Clarke-Hess 748 function generator and by co mparing, using the Tektroni x oscilloscope, the original applied waveform and the Vcpd signal detected by the system (Fig. 4.2). The period of the applied waveform was 10 ms while its am plitude was 220 mV (the amplitude of the signal reported in Fig. 4. 2 is 100 times attenuated). Figure 4.2. Comparison between the Vcpd signaldetected by the CPD system and the original voltage signal as observed via osci lloscope.This result confirms the possibility to measure fast (sub-ms range) tran sients with our CPD apparatus. As is evident from Fig. 4.2, the CPD system was able to perfectly detect the square waveform applied to the metal chuck. Rise and fall times of the detected signal were below 15 s. Stray capacitance. During preliminary measurements performed with the CPD apparatus we observed the capacitive couplin g of metallic cable connectors to the Monroe probe plates. After this finding, most of the metallic connecti ons in the apparatus were properly insulated to reduce stray capacitances.

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100 Sample charging operated by the vibrating Monroe probe. Leaving a semiconducting sample under the CPD probe for long times may cause an evident charging effect on the semiconductor surface. This charging effect is probably generated by the chemical species deposited by the vibrating shutters of the probe on the semiconductor surface. Hence, in all the experiments described in this and the next chapter, samples were never left under the vi brating probe for long time spans. For the case of long recovery times (e.g., typically required by hexagonal SiC) samples were removed from the apparatus and repositioned later. If vibrating probes are to be used for time-continuous CPD measurements we st rongly recommend the active use of the nitrogen purge insert which many of them in corporate. A plot of a significant charging effect caused by the probe on a 3C-SiC sample is reported in Fig. 4.3. Figure 4.3. Probe charging effect on the surface of a 3C-SiC epilayer within the first 18 hours. Note the decay in the observed CPD voltage caused by continuous measurement over the sample surface. Instrument performance verification. As a final step in th e calibration procedure, the instrument performance was compared to that of a similar commercial CPD apparatus from SDI. A comparison between the readin gs obtained at USF and SDI for 4H-SiC

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101 epitaxial samples are listed in Table 4.2. Considering th at surface potentials may vary significantly (up to tens of millivolts) on different areas of th e same semiconductor because of small topographical differences, cr ystal defects and impurities the results reported in Table 4.2 can be considered highly satisfactory. In conclusion, the USF CPD apparatus used in this work was well calibrated and, therefore, measurements made with this tool are reliable. Table 4.2 Comparison of CPD measurements of 4H-SiC epilayers using two CPD tools. Values obtained with the USF CPD sy stemare compared to those obtained with a similar system at SDI. SDI CPDM apparatus USF CPDM apparatus Sample ID Doping type s (mV) s (mV) USF-b-003 n -330 -340 USF-b-005 p 380 370 4.1.3. Procedure for CPD measurements of semiconductors in air ambient A specific procedure was developed to perform CPD measurements of semiconductors in air ambient. The procedure, described below, was used to perform all the CPD measurements that were used to ch aracterize the electronic behavior in ‘steady state’ and upon charging of the samples selected for the semiconductor-cell-electrolyte measurements. 1) Sample positioning over the vacuum chuck 2) Vacuum pump activation 3) Probe lowering at 1 mm from the sample surface

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102 4) Faraday cage closure 5) Measurement in dark Vcpd,dark 6) Switch on UV light 7) 10 second wait 8) Measurement in light Vcpd,light 9) Switch off UV light 10) Calculation of surface potential: s = Vcpd,dark Vcpd,light 11) Repetition of steps (5) to (10) for a second time In step (1) we always used metallic tweezers because charging effects were observed when handling the samples with Teflon or plastic tweezers. Since the Vcpd,light value was reached with different time intervals for different samples and it mostly depends on the quality of the crystal material (e.g., trapping effects may increase the time necessary to get to Vcpd,light), we selected the 10 se cond time interval between measurements after preliminary experiments indicated it as the ideal illumination time. The surface potential values reported in the next sections are expressed as mean ( S) standard deviation ( ) and are calculated from a tw o measurements distribution. 4.2. SiC and Si for cell-semiconductor inte raction studies: sample selection and description The final objective of this work is to study the electr onic interactions between biological cells and SiC samples via CPD. In order to perform successful CPD measurements of the complex semiconductor-ce ll-electrolyte system the CPD apparatus

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103 has been properly assembled and calibrated (§ 4.1). In this section we first discuss how we selected suitable samples for cell-semic onductor CPD investigations (§ 4.2.1) and then describe their characteristics (§ 4.2.2). Besides SiC samples, Si samples were also selected for cell-semiconductor interaction studies and are ther efore electronically characterized via CPD in this chapter. Si samp les are mostly used as controls to ensure the validity of the performed CPD measurem ents, since the surface potential of this material in response to different chemical processes has been characterized by past studies [112-114]. 4.2.1. Sample requirements for dry and wet CPD measurements The choice of suitable samples for the cell-semiconductor elec tronic interaction studies we aim to perform is of primary impor tance. In this secti on we describe which sample characteristics may negatively affect the dry (sample in air) and wet (sample in liquid) CPD measurements and how they influenced our sample selection. Doping concentration. High doping concentration in a semiconductor reduces the possible band bending range and, therefore, th e sensing potentiality when a charge is deposited on its surface. Therefore, we chose to work only with samples that presented doping concentrations NA,D < 1017 atoms/ cm3. Epilayer thickness. The penetration depth of the UV light used in our apparatus to generate the SiC fl atband condition (h = 3.4 eV) is approximately 8 m for the cubic and 11 m for the hexagonal poly type [115]. If the epilayer th ickness is lower than these values the possibility of ha ving interference in the CPD measurements caused by the generation of EHPs in the substrate underlyi ng the epilayer is significant. We found that for hexagonal p-type SiC epilayers grown on n+ bulk crystals the presence of a p-n

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104 junction relatively close (typical ly < 20 m) to the sample su rface affected the dark-light CPD measurements by causing Vcpd,light oscillation and a typical n-type behavior which was probably generated by majority carrier inj ection from the substrate into the epilayer. To avoid this problem, all the p-type samples selected for the CPD studies reported in the rest of this work were bulk crystals. On th e other hand, the presence and depth of SiC-Si heterojunctions in the 3C-SiC epilayers grown on Si seemed not to constitute a problem for CPD measurements. A specific experiment wa s designed to evaluate the effect of the heterojunction depth on the Vcpd readings: 3C-SiC epilayers whose Si substrates were removed and a series of 3C-SiC epilayers on Si with thickness va rying from 3 to 20 m were chemically charged and their surface potentials then measured. No significant differences were found among the magnitudes of their surface barriers which indicated an irrelevant effect of the heterojunction on th e final CPD readings. Therefore the 3C-SiC epilayers used in the rest of this work present variable thickness. Sample dimensions. In several instances we noticed that measurements of samples with areas smaller than 1 cm2 were less accurate and repeatable than measurements of samples with larger areas. This result can be easily justified by the existence of fringing effects arising from the fact that the samp le dimensions are close to the measuring electrode dimensions. Therefore, sa mples with areas larger than 1 cm2 were used for all the experiments described in this chapter. 4.2.2. Selected SiC and Si samples for cell-semiconductor CPD investigations The samples selected to perform the ce ll-semiconductor electr onic interaction studies, and which therefore will be electronicall y characterized in the next sections with different chemical charging procedures, were as follows: 3 off-axis n-type 6H-SiC(0001)

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105 epilayers grown on 6H-SiC and doped in the low 1016 cm-3 range; 3.5 off-axis pure grade p-type bulk 6H-SiC(0001) doped in the low 1015 cm-3 range; 3C-SiC(001) epilayers grown on Si(001) and unintenti onally doped n-type in the 1015 cm-3 range; n-type bulk Si(001) doped in the low 1015 cm-3 range; p-type bulk Si(111) doped 21015 cm-3. The samples were at least five for each category reported. All the samples had dimensions of at least 4.9 cm2 (e.g., a quarter of a two inch wafer) and presented extremely flat surfaces (see Fig. 4.4). In particular, all the SiC su rfaces presented atomic steps: the ones that originally presented polishing scratches were H-etched with processes suitable for the specific polytype (Chapter 2) to generate atomically flat surfaces and subsequently reduce the surface state density. Figure 4.4. AFM micrographs (22 m s cans taken in tapping mode) reporting the morphologies of the samples selected for CP D measurements: (a) ntype Si(111), (b) 3CSiC(001), (c) n-type 6H-SiC(0001). P-type Si(111) and 6H-SiC(0001) (not shown) presented surface morphologies simila r to (a) and (c), respectively. 4.3. Surface potential of SiC and Si in ‘steady state’ As we already pointed out in the introductory section, it is of primary importance to electronically characterize, via CPD, th e samples selected for cell-semiconductor electronic interaction studies. If the pattern of the typical electronic behavior of these

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106 samples in response to different chemical ch arging processes is defi ned, it should be easier to understand the more complex electroni c response that they may have in relation to electrolytes or cells. It is therefore of primary importance to define the initial band bending of these surfaces before any kind of surface charging or hydrogen etching process is applied. Since, in many instances after chemical treatment, the initial states (i.e. band bending) were also the ones to wh ich each specific sample had the tendency to come back to, we defined them as ‘steady states’. The samples analyzed in this section and in the rest of the chapter reached their ‘steady state’ after being treated with RCA cleaning in the past. To facilitate the reader in the analysis of the CPD results reported further in the text we recall that when a surface is charged positively the bands bend down ( s > 0) and when is charged negatively the bands bend up ( s < 0). Hence, for an n-type semiconductor the surface potential s is zero for flatband, ne gative for depletion and positive for accumulation. The opposite is true for p-type. Table 4.3 contains an elucidation of these concepts. Table 4.3 Semiconductor surface potentia l versus the overall surface charge. Surface charge > 0 Surface charge < 0 n-type s > 0, accumulation s < 0, depletion p-type s > 0, depletion s < 0, accumulation The ‘steady state’ conditions for the samples studied were the following: accumulation for p-type Si(111); depletion for ntype Si(111) as well as for pand n-type SiC(001). The observed conditions were in agre ement with what has been reported in the

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107 literature [114, 116] and suggest that all the samples were depleted of their majority carriers at the surface with the exception of p-type Si(111). The results obtained for Si and n-type SiC may be explai ned by the predominance of filled acceptor states near the surface. Instead, we suggest the presence of unoccupied donor states on p-type SiC surfaces. Fig. 4.5 is a band diagram representa tion of the ‘steady st ate’ condition for both Si and SiC surfaces. Obviously, the fact that the surface appeared to be negatively (positively) charged for na nd p-type Si and for n-type Si C (p-type SiC) does not mean that only filled acceptor (unfilled donor) states were present on the surface of these semiconductors, but that probably thei r effect was the predominant one. Figure 4.5. Band diagram representation of the ‘Steady state’ condition for nand p-type (a) Si and (b) SiC surfaces, respectively.

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108 4.4. Effect of hydrogen etching on the surface potential of SiC surfaces Besides producing atomically flat surfaces, H-etching should also terminate most of the dangling bonds of SiC surfaces with H atoms [117, 118], which should cause significant changes in the measured value of s. Moreover, if a complete chemical passivation of the SiC surface is achieved with H-atoms, th e resulting measured surface potential should be null [119]. From an elect ronic point of view this happens because hydrogenation replaces the surf ace states with Si-H bond ing and antibonding states, which are positioned below the valence-band maximum and above the conduction band minimum, respectively. Hence, no charge is transferred, which results in flatband ( s=0) [117, 118]. Therefore H-etching represents an additional technique, besides wet chemical treatment (§ 4.5), that can be used to modi fy the electronic properties of a surface and therefore to define its elec tronic behavioral pattern. The H-etching processes used to modify the surface potentials of 3Cand 6H -SiC samples were similar to the ones specifically developed in Chapter 2 for these polytypes, with the only difference that the etching times were reduced in a way that would ensure maximum H-termination but minimum surface morphology modification. Since all the SiC samples, selected for this experiment from the groups reported in § 4.2. 2, were already atomically flat at the moment of etching we did not want to ri sk surface over-etching which would increase surface state density. In fact, we are only intere sted in the chemical effect that processes have on the electronic state of the selected surfaces, and not on the influence of other variables such as su rface morphology. We therefore redu ced the etching time of 3C-SiC to 10 minutes and of 6H-SiC to 5 minutes.

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109 After H-etching, we found that 3C-SiC(001) was in a flatband condition. This is a promising result because it shows the possibility of a complete electronic passivation of 3C-SiC surfaces which was never reported be fore. Complete passivation of hexagonal SiC surfaces was reported by Seyller for n-type 6H-SiC(000 1) and p-type 6H-SiC(0001) [118]. In this section we first report the resu lts obtained by time-monitoring, via CPD, the H-etched 3C-SiC(001) surfaces (§ 4.4.1) a nd we then investigate, via XPS and ATRFTIR, the possible causes of the observed el ectronic passivation (§ 4.4.2). We also studied via CPD the effect of H-etching on n-type and p-t ype 6H-SiC(0001) (§ 4.4.3). However, in this case, electroni c passivation was not observed. 4.4.1. Electronic passivation by H-et ching of n-type 3C-SiC epilayers N-type 3C-SiC(001) epilayers were treated with 10 SLM of ultra-pure (grade 8.0) hydrogen for 10 minutes at 1200 C and AP. They were kept under hydrogen flow until a few minutes before their extr action from the CVD reactor, which took place at 400 C Just before the sample extraction the CVD reactor was purged with Ar to avoid dangerous combustion which otherwise may take place when opening the reactor door. The surface potential of th ese samples was measured before hydrogen treatment, immediately after and then monitored periodically over time. All the samples used in this experiment presented ‘steady state’ depletion values of s between -160 and -180 mV. All of them displayed, immediat ely after H-etchi ng, values of s between 0 and -7 mV which is indicative of complete electronic passivation. Fig. 4.6 reports the surface potential monitoring for different 3C-SiC epilayers up to 1000 hour s after the etching process.

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110 Figure 4.6. Surface potential vs. time of several H-etched 3C-SiC(001) epilayers presenting similar characteristics. A qua si-null surface potential was observed immediately after etch ing on all the surfaces. It is evident that the surface potential changed quite rapidly during the first few days after H-etching. In average, after 10 days (i.e., 240 hours), the surface potentials of these samples were within 10 mV from their final values, which have larger magnitudes than the initially measured quasi-null s. We confirmed these re sults over a large number of samples. Other experiments were performed to study the effect of the final hydrogen cooling temperature on the magnitude and time variation of s. Besides 400 C, the final hydrogen cooling temperatures studied we re 550, 1000 and 1200 C. The basic etching process applied was identical to the one desc ribed above with the only difference that after the hydrogen flow was interrupted (at 550, 1000 and 1200 C, respectively) all the samples were cooled down in Ar until thei r extraction from the CVD reactor which was performed at 400 C. We found that the surf ace potential time dependence of samples cooled down in hydrogen to a final temperature of 550 C was completely similar to what

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111 was reported in Fig. 4.6. In fact, immediately after H-etch ing, those samples displayed quasi-null s values and reached, in roughly 10 days, the final s value. On the other hand, although all the samples cooled down in hydrogen to higher final temperatures presented quasi-null s values immediately after H-etching, they very quickly abandoned this condition. The ones cooled under hydrogen to 1000 C displayed surface potential values of roughly 20 mV after 2 hours of air ex posure (Fig. 4.7, filled squares). In a more dramatic way, samples cooled down under hydrogen only to 1200 C presented surface potentials of roughly 20 mV after only 20 mi nutes in air ambien t (Fig. 4.7, filled triangles). Figure 4.7. Surface potential vs. time of H-et ched 3C-SiC(001) epilayers with final hydrogen cooling temperatures of 400 C (unfilled square s), 550 C (filled diamonds), 1000 C (filled squares) and 1200 C (filled triangles).Note that the time axis is logarithmic. From these results we can conclude that the stability of the electronic passivation of 3C-SiC(001) epilayers strongly depends on th e final hydrogen exposure temperature. In another instance we also demonstrated that etching processes which cause drastic surface

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112 morphology variations will not l ead to electronic passivation. Specifically, an atomically flat 3C-SiC epilayer treated with 10 SLM of hydrogen for 1 minute at low pressure and 1200 C and which presented, after the proces s, a much rougher and less ordered surface, did not display a quasi-nu ll surface potential. The effect of Ar annealing on 3C-SiC(001) epilayers was also investigated. Several samples with ‘steady state’ surface potentials of roughly -1 80 mV were Ar-annealed in the same CVD reactor used for H-etching at 1200 C and atmospheric pressure for 10 minutes. The results obtained are reported in Fig. 4.8. Figure 4.8. Surface potential time monitori ng, via CPD, of tw o Ar-annealed 3CSiC(001) samples with similar characteristics. Note that the time axis is logarithmic. It is evident from Fig. 4.8 that Ar-a nnealing did not provoke a quasi-null surface potential. However, the potential measured i mmediately after annea ling was significantly lower than the one presented by the same surfaces before the process and was totally comparable to the final s reported above for H-etched surfaces (which was in average 50 mV). We also observed that all the as-g rown 3C-SiC epilayers (e.g., never treated chemically after the growth process) displayed the same s values and time dependence

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113 than those reported in Fig. 4.8. This evidently indicates that the oxid e layer that forms as soon as the sample is extracted from the CVD reactor is very stable and provides a much better electronic passivation th an the oxide formed on the sa me surfaces after chemical treatment (e.g., after chemical treatment s < -100 mV for 3C-SiC always, as shown by Tables 4.6 and 4.7). This is also confirmed by the fact that some of the H-etched samples, which where HF dipped while they still presented a very low s, displayed, after HF treatment, less electronically passivated a nd more negatively charged surfaces (results reported in Table 4.4). Table 4.4 Charging effect of HF dip on H-et ched n-type 3C-SiC epilayers. Before HF dip the surfaces displayed a small | s| and an almost flatband condition. After HF dip the surfaces were evidently in a depleted condition. USF2-07-063b USF2-07-054.2 s before HF dip (mV) -29.6 -27.0 s after HF dip (mV) -140.9 -151.1 4.4.2. Characterization of passivated 3C-SiC surfaces via XPS and ATR-FTIR Since electronic passivation typically implies complete H-termination [118, 120], we decided to investigate the surface chemistry of the etched samples to probe if this was the case for our surfaces. Moreover, chemical studies can be used to motivate the differences in electronic behavior observed between H-etched and HF treated samples. The chemical characterization techniques used for this purpose were X-ray photoelectron spectroscopy (XPS) and atte nuated total reflectance Fourier transform infrared spectroscopy (ATR-FTIR). We performed XPS st udies, whose results are reported below, on an n-type 3C-SiC(001) sample before and after H-etching. The data obtained from the

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114 analysis performed before H-etching refe r to an HF dipped surface which presented s = -202.8 mV, while after H-etching the sample presented s = -7.1 mV. Surprisingly, a relatively high amount of oxygen was found on the etched surface (Fig. 4.9), which was comparable to the amount observed before etch ing (not reported). The magnitude of the O peak excludes the possibility that the oxygen observed is a minor contamination. Figure 4.9. XPS spectrum of a H-etched 3C -SiC(001) epilayer and relative elemental concentrations as calculated from the surv ey.A relatively high concentration of O was observed on the surface. However, the high resolution (HR) XPS scan s showed the existence of differences on the surface before and after etching. In particular, the graphiti c carbon concentration appeared to be lower for the etched surface (as is shown by the smaller shoulder in the HR C1s peak of the etched sample in Fig. 4.10(b)). Since it is quite easy to determine diffe rences in band be nding by photoelectron spectroscopy, we used the obtained HR spectra to confirm what was measured via CPD. Shifts in band bending can be measured vi a XPS because core level binding energies (e.g., Eb C1s) referenced to the Fermi level (EF) vary with surface band bending and

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115 because the sampling depth of XPS (2 nm) is much smaller than the width of the spacecharge layer. As is evident from Figure 4. 10(a) and (b) which report, respectively, the Si2p and C1s HR spectra obtained for etched and unetched 3C-SiC, the core level binding energies of the two samples ar e shifted by roughly 0.315 eV. The Si2p and C1s core level spectra of the H-etched sample shifted to higher binding energies, which is consistent with what was reported by Seyller [118] for n-t ype hydrogenated 6H-SiC(000 1) surfaces. From the energy band shifts observed with X PS we could deduce that the magnitude of the surface potential change between HF dipped and H-etched samples was around 300 mV, which is ~100 mV higher than the one m easured via CPD. This is possibly because the value was roughly calculat ed from XPS spectra which we re not performed with the aim of detecting band bending energy shifts. However, the shifts observed in the XPS spectra strongly confirm the existence of a different electronic behavior caused by different treatments on the same sample. Figure 4.10. Si2p and C1s core level spectra obtained for the same 3C-SiC epilayer before etching (bold line) and after et ching (light line).The core level spectra are shifted to higher binding energies after H-etching treat ment indicating a modification in surface potential.

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116 We should mention that even though the quasi-null surface potential measured for H-etched 3C-SiC epilayers may have strong ly suggested the presence of a complete chemical passivation operated by H atoms in a similar fashion to Si surfaces [120], the LEED data presented in § 2.2.3 for 3C-SiC epilayers treated with the same etching process provide evidence of a di fferent situation. In fact, if the etching process applied on 3C-SiC had produced a completely H-termin ated surface, the re sulting LEED pattern should have been an unreconstructed (11), wh ere instead we observed a (51) (§ 2.2.3). To better understand the cause s of the peculiar electronic behavior observed for the etched epilayers, we performed ATR-FTIR st udies which, differently than XPS, can provide interesting information regardi ng the presence and bonding of hydrogen on the surface. In these studies, performed under a constant nitrogen purge using a 45 beveled ZnSe crystal, we analyzed etched and HF dipped 3C-SiC samples with characteristics similar to those of the sample analyzed via XPS. The major differences that were observed by comparing the FTIR spectra of the etched and un-etched, HF treated, samples were the following: the spectra of th e etched sample exhibi ted peaks due to C-H bonds (Fig. 4.11), while the spectra of the un-etched sample presented a signal at wavenumbers corresponding to Si-OH bonds (F ig. 4.12). No peaks in the wavelength range of 2000-2100 cm-1 that could be associated to Si-H bonds were observed on either sample. However, the absence of Si-H peak s from the FTIR spectra of the H-etched sample, where they would be mostly expected does not imply that Si-H bonds were not present on the sample surface. In fact, si nce the ATR-FTIR apparatus used was not equipped with a polarizer, we could not perf orm p-polarized analysis, which can reveal bonds perpendicularly oriented to the surf ace as Si-H on SiC surfaces may be [121].

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117 Figure 4.11. ATR-FTIR spectra of H-etched 3C-SiC in a C-H stretch region indicating the existence of different typologies of C-H bonds. The major peak frequencies of the H-etch ed 3C-SiC epilayer in the C-H stretch region were located at 2849, 2917 and 2952 cm-1 and assigned to CH3 asymmetric, CH2 asymmetric, and CH3 symmetric bonds, respectively (F ig. 4.11). However, we also observed, by deconvoluting the spectra, sp3CH and CH2 symmetric stretching modes. From a joint analysis of the XPS and FTIR data, it appears that the 3C epilayer surface contains, after H-etching, both oxygen and hyd rogen, the latter be ing directly bonded to C. Amy et al. reported that hydrogenation of 3C-S iC(001) surfaces with a (32) LEED pattern leads to passivati on of the topmost dangling bond s but creates and stabilizes others below the top surface. Once the hydrogenated surface was exposed to molecular oxygen they observed oxygen incorporation below the top surface without any loss in the H coverage [120]. Even though our 3C-SiC( 001) surfaces presented a different surface reconstruction (i.e. (51)), it is still li kely that, upon molecular oxygen exposure, O atoms were incorporated below the top su rface without affecting the H-terminated topmost dangling bonds. Further studies are needed to investigate these interesting findings. However, the fact that H-etched 3C-SiC presented a high density of C-H

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118 bonding, sustains the suggestion made in § 2.2. 3 that the H-etched 3C-SiC epilayers may be C-terminated underneath the oxidic layer. Figure 4.12. ATR-FTIR spectra of un-etched HF treated 3C-SiC in a Si-OH stretch region displaying the ex istence of Si-OH bonds. The Si-OH peak observed for the HF dipped samples (Fig. 4.12) confirms that this chemical treatment leaves the SiC surface essentially covered wi th -OH groups [118]. The hydroxylic groups are probably th e cause of the depletion (i.e., s < 0) observed on the n-type surfaces after HF dip. Summarizing, the results presented in the last two sections indicate that the 3C-SiC surfaces are, after an appropria te H-etching process, electron ically passivated. However, at a chemical level the surface does not pres ent only hydrogen atoms, which is the case for Si [120], but also significant amounts of oxygen. The electronic passivation tends to deteriorate over time and under ambient conditions. However, the final s value at which the H-etched epilayers tend to stabilize is much lower in magnitude than the one measured for the same surfaces after HF trea tment, which evidently tends to add negative charges to the sample surfaces.

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119 4.4.3. Surface potential of H-etched 6H-SiC In this section we describe the effect of H-etching on nand p-type 6H-SiC(0001) surfaces. H-etching was performed at 1550 C fo r 5 minutes at atmospheric pressure and with a hydrogen flow of 10 SLM. Before th e etching process both the n-and p-type samples were found to be in depletion, whic h is a typical condition for SiC surfaces (§ 4.3). After etching, however, n-type 6H-SiC (0001) became more depleted while p-type 6H-SiC(0001) was characterized by a strong accumulation. Hence, H-etching seems to charge the surface of 6H-SiC samples negatively by probably filling acceptor states. The surface potential observed after etching did not vary significantly within the first 1000 hours (Fig. 4.13), which indicates that, ev en if the surface of these hexagonal SiC samples was not electronically pass ivated, it was stable over time. Figure 4.13. Surface potential time monitoring, vi a CPD, of H-etched n-type (triangles) and p-type (circles) 6H-SiC surfaces. No te that the x axis is logarithmic.

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120 4.5. Effect of chemical treatm ents on SiC and Si substrates Different chemical treatments were used to study the electronic behavior of Si and SiC samples upon charging. Specifically, Si and SiC surfaces were treated with HF and potassium permanganate (KMnO4) solutions and the magnitude of the band bending generated by the specific treatment on these sa mples was evaluated. Since HF dip is the last cleaning step performed before cell deposi tion (§ 5.1.3), determination of its effect on the band energy level is significan t. The information obtained in this section is used to model the general surface potential behavior of the samples selected for the semiconductor-electrolyte/semiconductor-cell-el ectrolyte systems and as a possible aid for understanding the results presented in Chapte r 5. Si substrates were used not only to study the surface potential of smaller bandgap ma terials after chemical treatment but also to verify the validity of the obtained results since the electronic status of Si surfaces upon chemical treatment is partially alrea dy known [112-114]. § 4.5.1 describes the experimental procedure used for charging th e Si and SiC surfaces while § 4.5.2 reports and discusses the results. 4.5.1. Chemical charging experimental procedure The HF treatment was executed in the following manner: sample dip in a 50:1 HF:H2O solution for 1 minute; thoroughly rinse in DI water; nitrogen dry. To charge the surface by using potassium permanganate th e samples were immersed, after the oxide was removed, in a 2:3 KMnO4:H2O solution warmed at ~ 35 C for 10 minutes, rinsed in DI water, and nitrogen dried. The char ging processes and the subsequent CPD measurements, whose results are reported below, were repeated for at least three samples of each typology listed in § 4.2.2.

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121 4.5.2. Band bending operated by chemical charging of the surface: results and discussion Before chemical treatment all the samples were in the ‘steady state’ condition described in § 4.3. Upon HF treatment, the p-type Si surfaces abandoned the accumulation condition and became depleted while the depletion of the n-type Si surfaces became less significant (Table 4.5). These results definitely show that HF treatment adds positive charges on Si surfaces which agrees with what was reported in [114]. This also indicates that the surface passi vation is not complete and that even if the surface is very likely largely H-terminated (which is conf irmed by the higher hydrophobicity of these surfaces), oxygen is still present (as confirme d by XPS analysis). However, it has to be mentioned that the significant s reduction observed for n-type Si(111) is likely an effect of the better passivation ope rated by a partial H-termin ation of the surface. Upon HF treatment, n-type and p-type SiC samples were still found to be depleted as they were before. However, as can be seen in Table 4.5, n-type SiC showed a greater depletion (e.g., more negative s value larger depletion width) while p-type SiC tended towards a less depleted state. This is indicative of a higher negative charge on SiC surfaces after the HF di p. It is known that HF treatmen t on SiC surfaces does not produce the passivating results that it does for Si: th in native oxides (~10 ) are typically present on the surfaces of ea ch SiC polytype immediately afte r HF exposure [116]. This was confirmed by the ATR-FTIR and XPS analysis performed on HF tr eated 3C-SiC samples and whose results are reported in § 4.4.2. The presence of hydroxilic groups -OH, detected on 3C-SiC surfaces via ATR-FTIR (F ig. 4.13), is a likely explanation for the increase in negative charge obs erved after HF treatment.

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122 In conclusion, all of the Si C and Si samples displayed different degrees of surface depletion after an HF dip: th e overall effect at the surf ace was probably the one of occupied acceptor states for n-type sample s and of unoccupied donor states for p-type samples. Table 4.5 Effect of HF dip on Si and Si C surface potential as measured via CPD. Sample Si n-type Si p-type 3C-SiC n-type 6H-SiC n-type 6H-SiC p-type ( S) ( ) (mV) before HF -219.9 0.6-60.3 2.8 -119.5 1.1-207.8 3.2 250.1 7.1 ( S) ( ) (mV) after HF -78.2 2.5 114.9 5.9 -168.3 0.1-240.7 1.9 108.3 3.8 Upon potassium permanganate treatment, whic h was performed after an HF dip, the Si(111) surfaces appeared to have acquired a more negative charge: n-type surfaces showed a greater depletion while p-type surfaces displayed a decrease in depletion (Table 4.6). On the other hand, the SiC response to potassium permanganate was conflicting: both n-type and p-type surfaces tended to a more depleted state (Table 4.6). Table 4.6 Effect of potassium permanga nate on Si and SiC surface potential as measured via CPD. Sample Si n-type Si p-type 3C-SiC n-type 6H-SiC n-type 6H-SiC p-type ( S) ( ) (mV) before KMnO4 -78.2 2.5 114.9 5.9 -300.9 7.1-339.9 11.8 189.5 14.8 ( S) ( ) (mV) after KMnO4 -188 4 70.3 0.9 -345 12.8 -1405.6 257.3 13.8

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123 Because of the oxidizing nature of potassi um permanganate, an oxide is present on both Si and SiC after the treat ment. This oxide introduces additional surface states, and therefore additional charge, Qit, at the interface with the semiconductors. Thus, the surface potential measured after permanganate treatment is probably largely influenced by the charge associated with Qit. Also in this case we can conclude that filled acceptor states were predominant at the oxide-semiconduc tor interface of n-t ype Si and SiC while unoccupied donor states defined the res ponse of p-type Si and SiC surfaces. Although the results reported above could be repeated at different times, it has to be considered that the effect of chemical char ging on both Si and SiC samples is extremely dependent on the topographical and electronic characteristics of the particular substrate used. It is possible that two samples presen ting the same doping and orientation may still have a different response to a specific chemi cal treatment. Still, the results presented were confirmed over a sufficiently large sele ction of samples (especially for 3C-SiC), which makes them statistically relevant. Tabl e 4.7 summarizes what is reported in this section. Table 4.7 Summary of the effect of HF and KMnO4on Si and SiC surfaces: charge added by the chemical treatment with respect to the initial state and sign of the surface potential measured. Si n-type Si p-type 3C-SiC n-type 6H-SiC n-type 6H-SiC p-type HF Add + charge. s < 0. depl Add + charge. s > 0. depl Add charge. s < 0. depl Add charge. s < 0. depl Add charge s > 0. depl KMnO4 Add charge. s < 0. depl Add charge. s > 0. depl Add charge. s < 0. depl Add charge. s < 0. depl Add + charge. s > 0. depl

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124 Both the HF treatment, which is performed before cell deposition, and the potassium permanganate treatment, which help s to reproduce the case when an oxide is present (e.g., immersion in cell culture me dia), depleted the Si and SiC surfaces of majority carriers. The presence of a depletion layer is idea l for sensing charges added on the surface and therefore is a good star ting surface condition for performing semiconductor-cell-electrolyte measurements. In fact, the addition of a fixed amount of charge on the semiconductor surface would cau se a much larger signal variation in surface potential when starting from a deplet ion condition as opposed to accumulation of the semiconductor. 4.6. Summary This chapter forms the basis for the semiconductor-cell-electrolyte CPD investigations that will be presented next in Chapter 5. The CPD apparatus has been properly assembled and calibrated while severa l preliminary measurements have allowed one to define the possible challenges that may be encountered when performing CPD measurements and to find suitable solu tions. Samples for the cell-semiconductor interaction studies have been selected, de scribed, and their electronic behavior upon chemical charging has been analyzed via CPD measurement. Thanks to the results presented in this and in previous chapters we are now ready to implement semiconductorcell-electrolyte CPD measurements for the investigation of electronic interactions between the biological world and semiconducting materials.

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125 Chapter 5. CPD Studies of th e Semiconductor-Cell-El ectrolyte System At this stage of the research SiC crystals with atomically-flat surfaces have been prepared and selected (Chapter 2, § 4.2. 2), their surfaces have been characterized morphologically (2.2.1, 2.3.1), crystallographically (2.2.3), chemically (2.2.3, 3.2), and electronically (4.3-5), while their biocompatib ility has been largely assessed (Chapter 3). Since all of the requirements li sted in Chapter 1 for the succ essful investigation of hybrid systems have been discussed and fulfilled, we can now describe the results obtained in the final part of this research. Understanding how the presence of charge s on the surface of a biological cell may affect the electronic band bending in a semiconductor would open a wide range of possibilities in the bio-sensing application area and is therefore one of the primary objectives of this work. In fact, up to da te, even though the possibility of electronic communication between electrogenically ac tive cells and semic onductors has been proved by cell electrical r ecordings operated by transist ors [122-125], the mechanisms underlying the electronic communication have not been studied. In this chapter we report the methodology and results that we used in the attempt to describe th e effect of the cell charge on the energy bands of a semiconductor. For this purpose, we selected dark/light non-contacting CPD measurement as a suitable technique capable of detecting the band bending at a semiconductor surface without di scharging it (§ 1.4) However, as we already pointed out in the introductory secti on of Chapter 4, the necessary presence of

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126 liquid in the CPD measurements we ai m to perform highly complicates these measurements. Nonetheless, thanks to the accurate calibration of our CPD system, we report how we were able to obtain vali d and repeatable CPD readings from semiconductor surfaces immersed in an electro lyte. Unfortunately, the results obtained from the CPD investigations of cell-semic onductor systems reported in this chapter did not display any measurable influence of the cell charge on the semiconductor band bending. However, the models and explanations used to describe th e obtained results are of primary importance for the future impl ementation of successful techniques for the investigations of cell-semiconductor electronic interactions. Indeed in Chapter 6 we will discuss the possibilities to continue this work, based largely on the knowledge gained from the results presented hereafter. Specifically, in this section we first descri be the experimental procedures developed for performing CPD measurements of semic onductor-cell-electrolyte systems (§ 5.1). Subsequently, we report the CPD characteriza tion results obtained during investigations of semiconductor-electrolyte interactions (§ 5.2). In this section, in addition to SiC, which was selected at the beginning of this research as the ideal substrate material for our CPD studies (§ 1.2), Si substrates were also us ed (i.e., the ones selected and electronically characterized in Chapter 4). Th e results presented confirm that SiC is preferred over Si for the CPD measurements we intend to perf orm. In § 5.3 we explain the reasons that brought us to choose the cell lines used in this chapter for cellsemiconductor electronic interaction studies. Subsequently, in § 5.4 we present the results obtained studying, via CPD, the effect of adherent cells on semi conductor surfaces. Interestingly, no electronic effect operated by the cells on the energy ba nds of the semiconductor was observed. For

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127 this reason we modified the approach and us ed, in the CPD investigations, non-adherent cells that are well known for their elec tronic charge: red blood cells (RBC or erythrocytes). In this instance, an eviden t cell-concentration-dependent band bending was observed on SiC surfaces (§ 5.5). Unfortunately, the apparent positive result obtained using RBCs was most likely caused by an optical effect associated with the nature of the erythrocytes rather than by their electroni c charge. Theoretical models and possible explanations for the results reported in secti ons 5.2, 5.4 and 5.5 are presented in § 5.6. 5.1. Experimental procedure for CPD measurements of the semiconductor-cellelectrolyte system There are several steps that have to be performed to prepare the complex semiconductor-cell-electrolyte system which is going to be measured via CPD. First, each selected semiconducting sample has to be cleaned with an optimized procedure which eliminates any organic residue from its surface and which leaves the semiconductor in a depleted condition (§ 5.1.1) Second, the sample has to be processed in a way that allows one to culture/deposit cells on its surface and that, at the same time, ensures good electrical contact of the sample backside with the metallic chuck (§ 5.1.2). Third, cells have to be seeded/cultured on the semiconductor surface (§ 5.1.3). Finally, the CPD measurements can be performed (§ 5. 1.4). In this section we describe all the procedures developed for each part of this experiment. 5.1.1. Chemical preparation Each sample that was used for cell-sem iconductor interaction st udies presented bioresidue on its surface at the end of the experi ment. Since throughout this work we re-use

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128 the same samples over and over to ensure result repeatability and to lower experimental cost, it is of fundamental importance to us e a cleaning technique which ensures complete bio-residue removal. Moreover, the chemistry used in the cleaning procedure has to leave the semiconductor in a depleted condition since a depleted electronic status increases the sensing potentiality of the semiconductor. The cleaning procedure that was selected for this purpose was Piranha followed by HF. In fact as we demonstrated in § 3.3, Piranha is the only chemical procedure that was found e ffective in complete bio-residue removal. The HF dip is used as a final step to ensure oxide removal (§ 3.1.1) and to generate a depletion region in the semiconductor surf ace (§ 4.5.2). Specifically, the applied procedure was constituted by the following steps: immersion in Piranha (2:1 H2SO4:H2O2) for 5 minutes; de-ionized (DI) water ri nse; dip in hydrofluoric acid solution (50:1 H2O:HF) for 2 minutes; thorough DI water ri nse. Sterilization of the samples via ethanol dip was considered unnecessary for two main reasons: 1) in many experiments cells were only deposited on the semic onductor surface and not cultured, the CPD measurements being immediately performed; 2) the strength of the Piranha and HF cleaning applied before cell de position were found to be e nough to avert the danger of bacterial contamination even for the longest culturing time which was fixed at 24 hours. It should be mentioned that in all of the experime nts performed and that will be reported in the next sections we never en countered contamination problems. 5.1.2. Sample processing In the final CPD measurement that we aim to perform a Monroe probe will be lowered on a semiconductor (typically SiC) whose surface, wet by an electrolyte, presents deposited/cultured cells and whose backside forms a quasi-ohmic contact with

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129 the metallic chuck of the CPD apparatus (§ 1.4.2). Therefore, the sample should be processed to ensure good electrical backsi de contact and a clean platform for cell culture/deposition. Obviously, since cells su rvive in aqueous environments and since aqueous environments oxidize semiconductors a first requirement is to physically separate the platform for cell culture (e.g., sa mple surface) from th e sample backside so that the electrolyte (e.g., cell media) will no t come into contact with it. Moreover, a way has to be found to avoid the sample backside oxidation which would surely take place for experiments where incubation of the samples in an atmosphere with 95% of humidity was required. For this purpose, prior to cell deposition, th e samples were processed in two different ways. This gave ri se to two different experiment al approaches that we have defined as the ‘free-standing’ sample appr oach and the ‘PEEK’ approach. In the free standing sample approach the liquid was confined in the top part of the sample by bordering the semiconductor edges with epoxy (i .e., crystal glue, see Fig. 5.1(a)) while Cu-Au contacts, evaporated on the backside of the sample, provided a stable electrical contact with the chuck. In the PEEK a pproach a machinable, autoclavable and biocompatible sample holder, made with Po lyetheretherketone (PEEK), was designed and fabricated. The backside of the sample was epoxied to the PEEK container while a hole with threads machined in the holder allo wed electrical back-side contact via a brass screw (Fig. 5.1(b)). The seal provided by the epoxy around the backside contact significantly slowed down the semiconducto r oxidation process that took place every time the sample and the sample holder were stored in the incubator for cell culturing.

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130 Figure 5.1. Sample configuration for cel l-CPD experiments showing (a) the freestanding sample approach and (b) the PEEK approach.In (a) the subdivision of the samples in parts with smaller area is represented. For experimental consistency cells and media (and media only for the experiments with electrolytes reported in § 5.2), were depos ited on a sample with an area of 4.9 cm2. For the PEEK experiments samples with this area were used while for the free-standing experiments samples with larger areas were partitioned using epoxy as described earlier (Fig. 5.1(a)). The two methods presented different a dvantages and drawbacks. In the freestanding sample approach the disadvantage was the toxicity of the epoxy which was in close contact with the cells while the a dvantage was quick sample processing. The drawbacks of the PEEK approach were l onger processing times and higher costs while the positive point was the biocompatibility of the sample holder. The morphology of B16 cells cultured on 3C-SiC samples processed in the two different ways is shown in Fig. 5.2. As expected, a higher quality cell mo rphology was found for the sample mounted on the PEEK sample holder (Fig. 5.2(a)). Cell s cultured on the fr ee-standing samples and located at sufficient distance from the e poxy displayed morphologies totally comparable

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131 to the ones of Fig. 5.2(a). However, we observed a poor morphology (i.e., minimization of the adhesion area) for cells cultured in th e vicinity of the epoxy, as expected (Fig. 5.2(b)). For this reason, CP D measurements of free-standing samples were performed far away from the epoxy, in areas where the cell morphology was equivalent to that observed on samples mounted in the PEEK holder. Figure 5.2. Cell morphology on (a) a sample mounted within the PEEK sample holder, (b) a free-standing sample in the vicinity of an epoxy drop (b).Note the different scale b ar in the two fluorescence mi croscopy images. In (b) an epoxy drop with a rounded are a was present in the centre resulting in non-adherent cells in this region. The choice of using two different expe rimental approaches eliminated the possibility that systematic errors in the a pproach chosen would affect the final results. Both the approaches were successfully used and yielded the same results as will be shown in the next sections. 5.1.3. Cell deposition / culture After the samples were clean ed and processed (e.g., mount ed in the PEEK container or bordered with epoxy), cells were either deposited (i.e., short-time experiments) or cultured (e.g., 24 hour experiments) on the semiconductor surface. The mammalian cells used in our studies were: B16-F10 mouse melanoma, HaCaT human keratinocytes and

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132 pig red blood cells (RBC or erythrocytes). While the fi rst two listed cell lines are epithelial cells that need to adhere to the s ubstrate to survive, RBC will not adhere to the SiC substrate. Therefore, two different expe rimental procedures were designed for the two cell categories. In the adherent cell expe riment cells were deposited on each sample at a density that allowed complete coverage of the sample within the first 24 hours. McCoy’s Modified Medium and Dulbecco’s Modified Eagle’s Medium (DMEM), both supplemented with 10% FBS, were the cult ure media used for the B16 and the HaCaT cells, respectively. Typical s eeding densities were of 10104 cells / cm2. After seeding, cells and samples (either epoxied to the PEEK container or free standing) were incubated at 37 C from a minimum of 4 hours (minim um time to ensure ce ll adhesion to the substrate) to a maximum of 24 hours in an atmosphere containing 5% CO2 and 95% relative humidity. In preliminary experiment s, CPD measurements were also performed after incubation times longer than 24 hours but they were fou nd to be highly unrepeatable and hence will not be reported. In the RBC e xperiments, erythrocytes were collected from pig blood after centrifugation at 200 g and subsequent plasma pellet extraction. Subsequently, they were re-suspended at a high concentration in phosphate buffered saline (PBS). The estimated density in PBS was 6.4109 cells/ml. In this case the CPD measurements were performed immediatel y after cell depositio n on the semiconductor surface. As mentioned in § 5.1.2, all the seedin g densities reported in this section refer to a sample area of 4.9 cm2.

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133 5.1.4. Experimental procedure for CPD measurements of semiconductor-cellelectrolyte systems The procedure used to measure the surface potential ( s) of a semiconductor immersed in an electrolyte and with adherent cells cultured on top was identical to the one reported in § 4.1.3. Steps (1)-(10) we re performed immediat ely after the sample (either free-standing or within the PEEK container) was re moved from the incubator and the media was reduced to a minimum amount. To ensure that differences in media pH (e.g., media becomes acidic over time because of cell byproducts) would not influence the surface potential of the sample under study, the culturing media was changed frequently and in particular 1 hour before th e execution of the CPD measurements. In this occasion the new media that was added did not contain fetal bovine serum (FBS) since it has been shown in electrophoret ic measurements that the pr esence of FBS may modify the charge associated with cells [25]. For RBC studies the free-standing appro ach was adopted because of its short processing time. The drawback of this approach which was represented by the toxicity of the crystal glue, was not an issue in this case since ce ll culturing on the semiconductor surface was not required and because of the ve ry short duration of the experiment. For CPD measurements of the semiconductor-RBC-el ectrolyte system a different procedure, which is reported below, was designed. 1rbc) Perform steps (1)-(10) reported in § 4.1.3 for the free-standing sample 2rbc) Add 100 L of PBS on the sample surface

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134 3rbc) Perform steps (1)-(10) reported in § 4. 1.3 for the sample-electrolyte system (taking care not to immerse the Monroe probe in the liquid) 4rbc) Add 1% of RBC concentrated solution 5rbc) Perform steps (1)-(10) reported in § 4.1.3 for the sample-RBC-electrolyte system (taking care not to immerse the Monroe probe in the liquid) 6rbc) Repeat steps (4rbc) and (5rbc) for higher concentrations of RBCs The surface potential values reported in th e next sections are expressed as mean ( S) and calculated from a two measurement distribution. The standard deviation values for measurements in liquid were similar to th ose of measurements in air, typically below 15 mV. 5.2. Electrolyte semiconductor systems Before studying the effect of cells on se miconductor band bending, we investigated the modification in surface potential that the presence of an electrolyte would cause on SiC and Si surfaces. The chemical behavior of Si in aqueous solution has already been investigated by several past studies [13, 14, 48, 102], which fac ilitates the understanding of the results that will be reported. All th e semiconductor-electrolyte interaction studies were performed using the free-standing sample approach, repeated several times (i.e., the results reported below within the tables are representative values), and depositing 100 L of liquid on the sample surface ( unless otherwise specified). Almost in every instance the electrolyte spread in a hydrophi lic fashion over the sample su rface. If not, a pipette tip was used to obtain a uniform coverage of liquid below the measuring electrode. In the

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135 presence of non-flat, high contact angle me dia drops on the surface (i.e., only obtained for one semiconductor-electro lyte combination, reported below), no valid CPD measurements could be performed. First, we evaluated the beha vior of n-type Si(111) and 3C-SiC(001) in water and in sucrose solution by adding 100 L of these liquids over the sample surfaces using a pipette. Before being in contact with the electrolyte, both the samples presented the typical depletion ‘stead y state’ described in § 4.3. CP D measurements were taken at several time intervals: immediately after the liquid deposition, afte r 5 minutes, after 10 minutes and finally after liquid removal. We observed a diffe rent time-dependent behavior for Si and Si C CPD values. In both cases aqueous solutions tended to reduce the initial depletion state, obtained after HF tr eatment, and caused the bands to shift towards a flat-band condition. On Si this was more ev ident being that the initial depletion region is thinner (e.g., in general smaller band bending amplitudes were observed for Si after HF treatment probably because of the better H-passivation, see § 4.5.2). However, we observed that longer times were required for the Si sample to reach its final s value while for SiC the transient was immediate (see Table 5.1). Also, after water removal, the s of SiC tended to come back towards the initial depletion value while Si remained in a basically flatband condition.

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136 Table 5.1 Effect of water on the surface pot ential of Si(111) and 3C-SiC(001) as measured via CPD. s (mV) of bare sample s(mV) after H2O at 0’’ s(mV) after H2O at 5’ s(mV) after H2O at 10’ s (mV) after H2O removal n-type Si(111) -135.5 -50.2 -35.1 -7.7 -10.1 n-type 3CSiC(001) -253.3 -117.2 -101.1 -108.3 -184 Even though Table 5.1 reports only the re sults obtained with water, completely similar results were obtained for the sucr ose solutions. The flatband condition observed on Si is probably justified by the smaller initial | s| value. However, the fact that Si surfaces did not tend to recover to the initial s value after liquid removal indicates that whatever caused the quasi-null observed s was still affecting th e dry surface. It is known that when Si is immersed in water a passivating oxide forms within the first 5-10 minutes [13]. Therefore, rather than an effect associated w ith the redox couples present in the liquid, the flatband condition observed is mo st probably related to the presence of the oxide that grows on the surf ace during the sample immersion in liquid. Very likely surface states which develop at the Si-oxide interface cause Fermi level pinning and yield to the observed behavior. The same quasi-null s was observed also when different electrolytes (e.g., culturing media, phosphate buffer saline (PBS)) were contacting the Si surface. Therefore, whether the final s observed for Si was due to Fermi level pinning or to a complete surface passivation, it is im possible to use Si as a substrate for cellsemiconductor electronic interaction studies. On the other hand, the existence of a still measurable depletion region on SiC surfaces contacting aqueous solutions encourag ed us to proceed with these studies. The

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137 electrolytes that were used in the experi ments with cells, namely McCoy, DMEM and PBS (§ 3.1.2), were deposited on the selected SiC samples with di fferent polytypism and different majority carrier types (e.g., nor p-type) (§ 4.2.2). CPD measurements of the semiconductor-electrolyte systems were take n a few seconds after the liquid came into contact with the semiconductor and repeated with in the first 10 minutes. Also in this case, no time dependent behavior was observed for the surface potential of SiC samples contacting electrolytes. T ypical results obtained fo r 100 L of deposited McCoy culturing media and PBS are reported in Tabl es 5.2 and 5.3, respectively. However, preliminary studies demonstrated that the am ount of liquid per surface area was irrelevant within reasonable concentrations (e.g., 50-150 L over an area of 4.9 cm2, see Table 5.4). The results obtained for DMEM culturing medi a were totally comparable to those in Table 5.2. Table 5.2 Effect of McCoy culturing media on surface potential of SiC and Si surfaces as measured via CPD. Values for Si reported for comparison purposes only. Sample Si(111) n-type Si(111) p-type 3C-SiC(001) n-type 6H-SiC(0001) n-type 6H-SiC(0001) p-type s (mV) of bare sample -228.4 113.1 -301.9 -337.6 287.7 s (mV) with 100 L of McCoy -16.1 26.1 -169.7 -131.4 180.8

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138 Table 5.3 Effect of phosphate buffer saline (PBS) on surface potential of SiC and Si surfaces as measured via CPD. Values for Si reported for comparison purposes only. not measured £ double values reported due to conflicti ng behavior in repeated measurements Sample Si(111) n-type Si(111) p-type 3C-SiC(001) n-type 6H-SiC(0001) n-type 6H-SiC(0001) p-type s (mV) of bare sample -129.8 100.5 -244.7 -354.2 222.7 / 214.2£ s (mV) with 100 L of PBS -14.9 0.8 -174.1 – 221.6 / 364.7£ Table 5.4 Effect of different amounts of McCoy culturing media on surface potential of 3C-SiC(001) as measured via CPD. s (mV) of bare sample s (mV) with 50 L of McCoy s (mV) with 100 L of McCoy s (mV) with 150 L of McCoy 3C-SiC(001) -301.9 -169.7 -165.7 -171 In Tables 5.2 and 5.3 surface potentials of nand p-type Si before and after electrolyte deposition are reported for compar ison purposes and also to confirm what was reported in Table 5.1. It is ev ident that the final quasi-null s value displayed by Si surfaces after electrolyte depos ition is independent of th e amplitude of the initial s value (e.g., even an initial s value as large as -228.4 mV appr oaches 0 mV when the surface is in contact with water). This indipendenc y strongly supports the hypothesis of a Fermi level pinning operated by surface states developing at the oxide-semiconductor interface. The presence of surface states at such a high density to cause Fermi level pinning on Si rather than on SiC surfaces c ould be due to the fact that Si samples were not prepared via H-etching and that they we re not flat at an atomic level (see § 4.2.2).

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139 It appears, that both McCoy and PBS tend to have the same effect of reducing the value of | s| (i.e., bands shifting away from deple tion) for all the SiC polytypes studied, with the exception of PBS on p-type 6H-SiC(0001). For n-type 6H-SiC(0001), s values could not be calculated because of the hydrophobic reaction that these substrates had with PBS: all the attempts to deposit any amount of PBS on the surfac e resulted in the formation of a high contact angle drop th at made it impossible to perform CPD measurements. On the other hand the result obta ined for p-type 6H-S iC were conflicting: either no effect on band bendi ng or a shift of bands toward depletion upon PBS exposure was observed. These results wi ll be modeled in § 5.5. 5.3. Cell line selection and properties For the cell-semiconductor electronic intera ction studies three different cell lines with different characteristics were chosen. Unfortunately, to date, no clear estimation of the cell charge of adherent mammalian cells has been reported, with the exception of RBCs. This is one of the main reasons why our investigations are truly pioneering and, at the same time, particularly challenging. The only technique that currently reports the evidence of the existence of a net cell surface charge is el ectrophoresis, which determines the mobility of cells in aqueous solution upon ap plication of an electr ic field between two electrodes. Thanks to these studies it is nowadays accepted that mammalian cells display a negative surface charge. Once th e electrophoretic mob ility is defined, it is possible to calculate the potential at the cell shear plane (i.e., zeta potential ( )) by using the von Smoluchowski equation = 4 /D where is the fluid viscosity, is the electrophoretic mobility of the particle (i.e., cell) and D is the dielectric constant of the fluid [6, 24]. Clearly, the value of by itself is meaningless if not associated with a pH and an

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140 electrolyte. The von Smoluchowski equation has been long debated in the past and presents major interpretative difficulties [126]. The first is sue in the adoption of this equation, for the determination of the zeta potential, resides in the fact that the relation was created for spherical particles whereas ce lls generally deviate from this assumption [6]. Also, this equation seems to be reliabl e only for zeta potentia l values < 120 mV and electrolytes containing more than 10-3 molar salt [127]. However, since it has been reported that erythrocytes made spherical by saponin have electric mobility, within the limit of experimental error, identical to the ones of the disc-shaped cells suspended in the same buffer, the von Smoluchowski equation has been largely adopted in the past for the determination of of erythrocytes [6, 24, 128]. Typical values of for erythrocytes vary between -15 mV and -30 mV, depending on the electrolyte of suspension [6, 24, 128, 129]. In general, for small values of the Debye approximation for the calculation of an erythrocyte charge density ( ) can be used and the relation between and is straightforward: = D /4 As is intuitive, high electr ophoretic mobilities lead to high zeta potentials and, therefore, are indicativ e of a high cell surface charge. Thanks to electrophoretic experiments and theoretical calculati ons performed using the aforementioned equations, the surface charge density of mammalian RBCs has been estimated to range between 63 and 188.810-8 C/cm2 [6, 128]. On the other hand, values of for other mammalian cells have not been calculated due to their incompatibility with the von Smoluchowski equation. Therefore, all knowledge about the possible char ge associated with a specifi c cell comes from chemical studies. From a chemical point of view, it is well known that the negative charge of mammalian cells is mainly caused by sialic acid, hyaluronic acid and chondroitin sulfate

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141 which are all present in the ex ternal part of the cellular membrane [130]. In particular, sialic acid is the only major constituent w ith negative charge in RBCs. Also, medical reviews in the past have reported results that tend to display higher electrophoretic mobilities for malignant cells [7]. The reduction of electrophoretic mobility observed for different malignant cell lines upon X-irradiation te nds to confirm this belief [5]. Known charge-related characteristics of the cells used in this work, together with a motivation for their use, are summarized below. 1) HaCaT human keratyinocytes. No known surface charge density. Estimated mobility comparable to those of healthy mammalian cells and therefore largely lower than erythrocytes mobility in saline solu tions [7]. Fast growing cell line, good for preliminary experiments. 2) B16-F10 mouse melanoma cells. No known surface charge density. Electrophoresis experiments performed for another B16 sub-line (B16-C2W) report mobility values comparable to those measured for human erythrocytes (e.g., = -1.05s1V-1cm) [5]. This could lead to the assumption that the is similar for the two cell lines. However, an approximately 7 times larger di ameter for the mouse melanoma cells (i.e., 40 m for B16 vs. 6 m for RBC) complicates the matter. Moreover, aggregation of B16 in clusters is often observed which may lead to the conclu sion of limited electrostatic cell-to-cell repulsion and therefore a lower than the one measured for RBCs. Malignant cell line, possible higher negative surface char ge with respect to HaCaT. Fast growing cell line. 3) Pig RBCs. Zeta potential and surface charge de nsity in saline estimated to be roughly -13 mV and 11.210-7 C/cm2, respectively [6, 128]. Po ssibility of obtaining

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142 significant electronic effect on the energy bands of SiC semiconductors. In fact, the total charge necessary to significantly deplete a 4H-SiC epilayer is on the order of 10-7-10-8 C/cm2. Depositing a monolayer of RBCs a bove the semiconductor surface would generate a uniform charge density well above the minimum value necessary to generate a measurable band bending on SiC (e.g. 11.210-7 C/cm2 > 10-7 C/cm2, which is required to completely deplete a 4H-SiC epilayer) [ 45]. Easy to obtain in large quantities. 5.4. CPD investigations of adhe rent mammalian cells-SiC systems §5.1 perfectly described the protocols used in the different stages that characterize CPD investigations of cell-semiconductor system s. In this section, before reporting the results obtained in the case of mammalian adherent cells, we will first describe the adopted experimental approach. In these e xperiments, when using the PEEK approach, we first evaluated the surface potentials of two identical sample s after the cleaning process ( s1,2 in Fig. 5.3), and then we probed whether the adopted sample processing technique (e.g., epoxy to PEEK in this case, § 5.1.2) had an effect on these values ( s1’,2’ in Fig. 5.3). We always observed that the sa mple processing technique did not modify the surface potential values measured before processing ( s1,2 = s1’,2’). Subsequently, B16F10 mouse melanoma and HaCaT cells were de posited, as described in § 5.1.3, on only one SiC substrate and cultured for different tim es ranging from 4 to 24 hours. The other sample was incubated only with culturing me dia (e.g., no cells) for the same amount of time. The CPD measurements were performed at regular intervals within this time range and the readings obtained for the substrate with cultured cells ( s1’’) were compared to the one obtained for the identical substrate (always cut from the same original wafer and treated in an identical fash ion) which was exposed to th e electrolyte but no cells ( s2’’).

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143 In particular, the difference between the su rface potentials measured when cells+media (or media only) were present and immediately before cell seeding (e.g ., after processing) was calculated for each sample ( 1,2 = s1’’,2’’– s1’,2’) and compared. Fig. 5.3 contains a flow chart describing the e xperimental procedure used when adopting the PEEK approach. For the free-standing approach the on ly difference was that typically half of a two inch sample (i.e., area A = 9.8 cm2) was partitioned in two subdivisions (similarly to Fig. 5.1(a) but only for a half wa fer) displaying an area of 4.9 cm2 each: one half was seeded with cells and media and the other on ly with media. Therefore, we used one sample divided in two parts instead of tw o identical samples. The latter approach eliminated the possibility that two diffe rent samples would undergo other surface modifications than the ones caused by cells when stored separately in the incubator.

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144 Figure 5.3. Flow chart descri bing the experimental procedur e adopted to investigate the effect of cell charge on semiconductor band bending. s1,1’,1’’ ( s2,2’,2’’) are the surface potentials for the sample with cells+media (media only); 1 ( 2) is the surface potential difference between the sample with cells+med ia (media only) and after processing. The procedure reported is for the PEEK approach. For the free-standing approach all the steps are the same only, instead of using two identical samples, we operate on two parts of the same sample partitioned by epoxy. Typical surface potential values obtained for B16 and HaCaT cells are reported in Tables 5.5 and 5.6, respectively. In Table 5.5 results are reported for 24 hour inspection of 3C-SiC(001) samples mounted on PEEK. However, they were also confirmed for n-

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145 and p-type 6H-SiC(0001), using the free-stan ding sample approach, and for different incubation times. The experiment was repeated at least 10 times yielding always the same results. The media used for this cell line was DMEM throughout the experiment. Fresh DMEM (without FBS) was replaced for both the samples, with seeded cells and without, one hour before the CPD measurements as desc ribed in § 5.1.4. The results reported in Tables 5.5 and 5.6 clearly show that no measur able differences in surface potentials were observed between the samples with cells+media and the samples with only media. Table 5.5 Non-detected electronic interact ion between HaCaT cells and 3C-SiC(001) after 24 hours from seeding (standard deviation < 15 mV). Sample ID s (mV) after cleaning + sample processing s (mV) of sample w/ media only s (mV) of sample w/ media and cells (mV) n-type 3C-SiC USF-06-271.1: media + cells -211.4 ( s1’) – -109.6 ( s1’’) 101.8 ( 1) n-type 3C-SiC USF-06-271.2: media only -207.6 ( s2’) -111.5 ( s2’’) – 96.1 ( 2) It appears evident from Table 5.5 that no significant differences were observed between the values calculated for the two samples. In fact, 1 – 2 = 5.7 mV which is within the standard deviation range for these measurements (e.g., 0 mV < < 15 mV, as reported in § 5.1.4). Table 5.6 illustrates the t ypical values obtained on di fferent polytypes of SiC processed with the free-standing sample approach after 4 hours from B16-F10 cells seeding. The same results were obtained also when using the PEEK approach and after different incubation times.

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146 Table 5.6 Non-detected electronic interact ion between B16-F10 cel ls and n-type 3CSiC, n-type 6H-SiC and p-type 6H-SiC 4 hours after seeding ( < 15 mV). s (mV) after cleaning + sample processing s (mV) of sample w/ media only s (mV) of sample w/ media and cells (mV) n-type 3C-SiC USF-07-039a: media + cells -276.1 ( s1’) – -175.4 ( s1’’) 100.7 ( 1) USF-07-039b: media only -284.4 ( s2’) -178 ( s2’’) – 106.4 ( 2) n-type 6H-SiC CO24a: media + cells -323.5 ( s1’) – -594.6 ( s1’’) -271.1 ( 1) n-type 6H-SiC CO24b: media only -349.8 ( s2’) -614 ( s2’’) – -264.2 ( 2) p-type 6H-SiC BQ02a: media + cells 215.6 ( s1’) – 252.6 ( s1’’) 37 ( 1) p-type 6H-SiC BQ02b: media only 232.2 ( s2’) 281.7 ( s2’’) – 49.5 ( 2) Also in this case, no statistically signif icant differences were observed between the s calculated for each pair of samples of the same polytype. Therefore, no electronic interaction of adherent B 16-F10 mouse melanoma cells with the SiC samples was observed. In every experiment, after the 24 hour CPD measurements were taken, all samples with seeded cells were inspected via fluorescence microscopy to evaluate cell morphology and coverage. We observed satisf ying morphologies and a coverage of at least one monolayer in all cases. The obtained results are justified and modeled in § 5.6.

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147 5.5. CPD investigations of RBC-SiC systems The CPD studies performed for B16-F10 melanoma and HaCaT cells on SiC surfaces did not display an evidence of band bending on the semiconductor surface. As we already mentioned in § 5.3, the surface char ge density associated with these adherent cells may not be sufficient to induce a measurab le effect on the elec tronic status of the semiconductor. On the other hand, the presen ce of a significant, and therefore easily measurable, electronic charge on RBCs and the fact that these cells are obtainable in large numbers brought us to choose them as an alternative, and hopefully more successful, cell line for the CPD experiments. We started the RBC experiments by depositi ng, on bare free-standing SiC samples already fully characterized via CPD, only th e electrolyte for our experiments (PBS), repeating what was done in § 5.2, only this time not for modeling purposes but to obtain a specific surface potential value for each sample to compare with the ones obtained after cell deposition. After this, we proceeded by adding gradually increasing densities of RBCs and by measuring, after each addition of cells, at least two su rface potential values to ensure statistical relevance of the data (as don e in the rest of this chapter, e.g. § 5.1.4). The amount of RBC concentrated solu tion (estimated cell density: 6.4109 cells/ml, (§ 5.1.3)) deposited via pipette each time were 1, 10, 20 and 50% of the PBS solution. The experiment was performed on at least three selected samples for each polytype and was repeated two times. The results obtained appare ntly showed a strong de pendence of the surface potential on the amount of cells present in so lution and, in all case s, the bands of the semiconductor appeared to shift away from depletion and towards flatband. The observed

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148 surface potential behavior is plotted in Figures 5.4, 5.5 and 5.6 for n-type 3C-SiC, n-type 6H-SiC and p-type 6H -SiC, respectively. Figure 5.4. Measured CPD surface potentia l for increasing amounts of RBCs for different 3C-SiC(001) samples. The surface pot ential values of the bare samples and o f the samples with PBS are also reported for comparison. As is evident from Fig. 5.4, the behavior of the band bending upon PBS deposition was consistent with what was reported in § 5.2. A significan t variation in the measured surface potential was observed even after addition of only 1% of RBC concentrated solution in PBS. However, this major effect co uld be easily justified by the fact that the total amount of cells deposited in this cas e was significantly high: approximately 6.4106 cells were present in the saline at the mo ment of the measurement. When adding an amount of cells equivale nt to roughly 3.2108 cells (50% RBC in Fig. 5.4) the calculated surface potential value approached zero.

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149 Figure 5.5. Measured CPD surface potential for increasing amounts of RBCs for two ntype 6H-SiC(0001) samples and data repeatab ility in two different experiments (DAY 1 and DAY 2) as calculated for one sample. The surface potential va lues of the bare samples are reported for comparison. Very similar results were observed for n-t ype 6H-SiC (Fig. 5.5). However, in this case, as explained in § 5.2, we could not monitor the effect of PBS on the electronic status of the semiconductor because of the high contact angle of the PBS drop with ntype 6H-SiC and the impossibility to flatte n it. However, once the RBCs were added in the solution, we observed a significant change in the electrolyte behavi or that, this time, flattened on the semiconductor su rface without resist ance (i.e., lower contact angle). This was probably due to a modification in the elec trostatic interactions between the n-type 6H-SiC surface and the saline buffer once RBCs were added, implying that RBCs were the main cause of the observed modification. In Fig. 5.5 the repeatability of the surface potential values obtained for the same n-type 6H-SiC sample in two different experiments can be also observed.

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150 Figure 5.6. Measured CPD surface potential for increasing amounts of RBCs for two ptype 6H-SiC(0001) samples.The surface potential values of the bare samples and of the samples with PBS are also reported for comparison. As is evident from Fig. 5.6, PBS alone te nded to shift the bands of p-type 6H-SiC towards depletion while even the smallest am ount of cells drastica lly changed the trend and apparently shifted the bands away from depletion and towards flatband. It has to be mentioned that the cause of the quasi-null surface potential observed for SiC samples in the presence of a high number of RBC’s (i.e., ~ 3.2108) was surely different than the one that generated the qua si-null surface potential on Si surfaces wetted by electrolytes (§ 5.2). In fact, for the Si C-RBC system, we observed that, immediately after PBS-RBC solution removal, the SiC energy bands shifted back towards the initial values that they presented prio r to liquid deposition: this i ndicates that the effect observed via CPD is strictly related to the presence of cells over the surface and not on absorbates or passivating layers forming on the semiconductor surface. The results reported were repeatable and al ways consistent. An example of the data repeatability is shown in Fig. 5. 7 for two samples of 3C-SiC(001).

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151 Figure 5.7. Repeatability of the surface potentia l values calculated for two different 3CSiC(001) samples from two differe nt experiments (DAY 1 and DAY2). Although appealing at first glance, the results obtained in the RBC-SiC CPD investigations presented some basic incongruencies. The most striking resides in the fact that RBCs are well known to display a negative total charge: the CPD readings obtained for p-type surfaces agree with this (shift towards accumulation indicates negative charge on the surface), while the results obtained fo r n-type surfaces suggests the presence of positive charges. Also, the fact that the su rface potential values calculated from the readings tended to zero both for p-type a nd n-type seemed somewhat suspicious. For these reasons, different hypothe ses were evaluated to yield satisfactor y answers to the observed conflicting behavior. As will be discussed shortly, the results reported in this section do not, unfortunately, describe the effect of cell charge on semiconductors, but are caused by optical issues induced by the erythrocytes.

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152 5.6. Discussion and modeling of the obtained results In this section we will discuss the results pr esented in § 5.2, 5.4 and 5.5 and we will also introduce and develop theoretical models which may help to better understand what was observed while performing CPD measuremen ts of the SiC-electrolyte-cell system. First, we will find reasonable explanations for the results obtained in § 5.2. In order to understand what is reported further in the text the reader should refer to § 1.4. As already discussed in Chapter 1, the semiconductor-electrol yte interface is extremely complex and can be modeled by using the el ectric double layer model. Maintaining the approximation of negligibility of the Helm holtz capacitance and of the surface state capacitance, which is reasona ble in the studied case (Csc << CH and Csc >> Css Ccpd = Csc, § 1.4.5), we can attribute the CPD voltages measured entirely to the excess of charge present in the space charge layer (e.g., Vcpd = Vsc). Recalling the basic principles from § 1.4.4, there are two diffuse layers present at the semiconductor/electrolyte interface: the Gouy layer in the liquid and the space charge la yer in the solid. On the solution side, ions distribute over the semiconductor surface at a minimum distance of ~ 3 (e.g., outer Helmholtz plane) as opposed to the excess char ge present in the space charge layer, and then diffuse within the Gouy layer until ch arge neutrality is reached at a distance OHP+WGouy. It has been estimated that the maximum distance at which charge equilibrium is reached within an aqueous so lution, and therefore where the potential drop entirely occurs, is roughly 30 nm [13]. Using the electrical double laye r model, the result obtained in § 5.2, where de pletion was observed in all instances upon electrolyte deposition on the semiconductor surface, can be schematically represented by an overall negative (positive) charge present on the ionic side for the depleted n-type (p-type).

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153 Schematic representations of the charge dist ribution in the electrical double layer for ntype SiC and of the associated potential varia tion in the electrolyte are reported in Fig. 5.8. Figure 5.8. Schematic illustrati on of the electronic status at the n-type SiC/electrolyte interface and relative potentia l variation in the fluid. From an energy band diagram perspective, the obtained results indicate that the redox energy levels (ERedOx) of the media used were below the Fermi levels of the n-type semiconductors and above the ones of p-type sa mples, which is likely (see Fig. 5.9). However, this explanation is not sufficient to describe the obtained results. This is because the Fermi level of nand p-type Si C samples typically lie only within 200 meV from the conduction band minimum (CBM) and the valence band maximum (VBM), respectively. Therefore, if we assume th at the redox energy level of the specific electrolyte lies in the middle of the SiC bandgap this would cause, once the electrolyte and the semiconductor are brought into contact a shift of the bands of roughly 1 V of magnitude towards depletion both for n-type and p-type samples. Th e energy band shifts reported in § 5.2 upon electrol yte deposition on the sample surfaces were of much

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154 smaller magnitude (~ 100 mV) and, most impo rtant, in almost all the instances (with exception of the p-type SiC-PB S combination) away from de pletion, which suggest that other factors, besides the positioning of ERedOx, influence the final CPD readings. Adsorbed ions and charges trapped by su rface states at the surface are possible explanations for the decrease in the positive excess charge in the space charge region that otherwise should be observed. Figure 5.9. Energy band diagram for n-type (L HS) and p-type (RHS) 6H-SiC/electrolyte interface assuming that the Fermi level pos ition in the semiconductor is 200 meV from the conduction and the valence band edges, respectively.Top: band diagrams before contact. Bottom: band di agrams after contact. Let us now develop a valid model to expl ain the results obtaine d in § 5.4. First consider the situation where only the electrolyte is presen t above the media which will cause a limited surface depletion like the one depicted in Fig. 5.8. Let us now perturb this situation by introdu cing a cell. As we already reported in § 5.3, the surface charge of mammalian cells is negative. The cell membrane is highly

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155 insulating [131, 132], which allows us to neglect the influence of the internal charge and to model the cell as a solid body with a charge density, concentrated on its surface. First we will consider the cell floating in the liquid, before cell adhesion to the substrate starts (e.g., time required for a mammalian cell of the lines used to adhere to a substrate is roughly 4 hours). Again, for a solid/liquid in terface (cell membrane/electrolyte) we will observe an electrical double layer with an a ssociated potential decreasing with increasing distance from the cell outer membrane. The pot ential at the cell shea r plane is known in the literature as the particle’s zeta potential ( ) as shown in Fig. 5.10. Figure 5.10. Schematic representation of a negatively charged cell suspended in liqui d and the system relative potential diagram. Note an electrical double layer forms on the cell surface. Next we will provide the best case es timation of the zeta potential and of the potential decay (e.g., best case = higher valu es), which will allow us to develop a model capable of explaining the results presented in § 5.3. For this purpose we will use the data

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156 found in the existent literature relative to red blood cells whic h, as explained in § 5.3, are most likely the cells with the higher charge density and therefor e the ones presenting higher zeta potentials and electros tatic effects in the outer regi on. Several studies agree in reporting that the zeta potential of RBCs in saline solution is approximately -15 mV and that this potential decays within the first 20 nm from the cell surface [6, 24, 129]. The negatively charged cell surface immersed in media and surrounded by counter-ions and the potential variation associated with it ar e schematically represented in Fig. 5.10. Let us now consider the case where the cell adheres to our semiconducting surface. As explained in Chapter 3, mammalian cells adhere to substrates via adhesion plaques which are approximately 15-20% of the cell contact area [131]. With an optimistic assumption we can approximate the contact area of a cell to be 50% of its whole area (see Fig. 5.11), which makes the adhesion plaque 7. 5-10% of the total cell area. Apart from the contacts, the average gap between the cell and the substr ate surface has been reported to be in the range of 50-150 nm [131]. Th e situation is schematized in Fig. 5.11. Figure 5.11. Schematic illustrating the limited electronic effect of the cell charge on the semiconductor. The Gouy layers for the cell -media (~ 20 nm) and for the semiconductormedia (30 nm) systems are indicated.

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157 It appears that, for the assumptions previously made and in the most optimistic case, only 10% of the total cel l charge has a direct effect on the semiconducting substrate. The remaining charge is too far away from the semiconductor surface to possibly have any effect on its electronic status. As explaine d in § 5.3, it is reasonable to assume that the charge associated with the adherent cells used is significantly lower than the one that erythrocytes display. This causes even lowe r zeta potentials and s horter potential decays which, combined with the reduced cell-su bstrate contact area, may lead to the nondetected electronic interaction between adherent mammalian cells and SiC surfaces. Last in this section we will discuss the re sults presented in § 5.5. As we already mentioned, the results obtained from CPD i nvestigations of RBC-SiC systems differ from the ones we would theoretically expect because the surface potentials observed for nand p-type samples upon addition of RBC did not be have in a fashion consistent with the addition of negative charges. Moreover, th e tendency of the meas ured potentials to converge to a null value upon addition of in creasing RBC concentrations both for nand p-type samples did suggest that the obt ained results were probably caused by a measurement-related deficiency rather than by the electri cal charge associated with RBCs. We identified two possi ble causes of the observed be havior: 1) a pH related decrease of the Helmholtz capacitance contri bution; 2) a decreased light penetration within the erythrocyte solution. Since the Helmoholtz capacitance (CH) is strongly dependent on the pH of the elect rolyte solution [13], a decrease in its value influenced by an increase in the pH, which could have been caused by the addition of erythrocytes in PBS, may have affected the final measured Vcpd by making the Helmholtz layer contribution non-negligible. However, this po ssibility was soon eliminated after the pH

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158 of erythrocytes was found to be extremely cl ose to the one of the PBS solution (e.g, 7.33 for the erythrocytes compared to 7.4 for PBS) [133]. Since the Helmho ltz potential varies by 60 mV per pH unit, such a small pH vari ation would not cause the significant effects observed in § 5.5. On the other hand, the reduction of light penetration within the erythrocyte solution immediately appeared as a valid explanation for the obtained results. In fact, it has been reported in optical studies of intravascular structures that the refractive index ( ) mismatch between erythrocyte cytoplas m and blood plasma causes strong light scattering [134, 135]. Specifically, the re fractive index of the erythrocyte is = 1.4 vs. = 1.337 presented by the blood plasma. The ma in component in the erythrocyte that influences the final refractive index of the cell is hemoglobin, which presents = 1.615 and makes up roughly 97% of the entire er ythrocyte dry content [136]. The same scattering mechanism exists for erythrocytes su spended in saline solution, which presents a refractive index equal to that of water ( = 1.33). Since all the experiments with RBCs were performed using PBS and since the ca lculation of the CPD voltage under deep illumination was made by shining light through the erythrocyte containing solution, it appears very likely that part of the light shone did not reach the semiconductor surface due to photon scattering by the RBCs. Moreover, another light related phenomenon, that probably has an even stronger effect on the fi nal CPD results reported in § 5.5, takes place upon illumination of erythrocytes: light abso rption. Also in this case, hemoglobin is mainly responsible, causing major absorption at lower wavelengths ( ), specifically for < 400 nm [137]. Since the wavelength of the LEDs used in our CPD system was = 370 nm, and the hemoglobin penetration depth at th ese wavelengths is at most 20.6 m, it is correct to assume that part of the light shone through the PBS-erythrocyte solution was

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159 absorbed and a portion, as aforementioned, was scattered. This combined mechanism is the most likely cause of the apparent decreas e in the surface potential s measured in § 5.5 upon addition of increasing amounts of RBCs. In fact, the higher th e concentration of RBCs in PBS solution, the more likely it is that a single photon will not reach the semiconductor surface because of scattering or absorption mechanisms. Hence, the low light intensity reaching the sample surface was probably not sufficient to create the flatband condition that characterized the m easurement upon strong illumination and the Vcpd,light value tended to match the Vcpd,dark value. Specifically, we observed a s reduction upon addition of 1% RBCs in PBS: in this case the cell density ( c) in the solution was 64103 cells/L. Ideally assuming that the cells are points equally spaced within the solution we can calculate a total of 3 c = 40 cells along one of the dimensions of a 1 mm3 cube. This assumption implies that the average distance between RBCs would be dc = 1 mm / 40 = 25 m. Considering their biconcave shape and assuming an average di ameter of 6 m and thickness of 1.2 m [138] their average distance is found to be 23.2 m. Because of the concave shape of the cells, the probability that the plane of inci dence and of exit for a single incident photon are parallel strongly de creases, which, applied for a larg e number of cells, leads to a significant scattering mechanism (Fig. 5.12) Furthermore, hemoglobin is a high percentage within the erythrocytes cytopl asm. Assuming a maximu m light penetration depth within the intracellular hemogl obin solution of 20.6 m and a minimum erythrocyte dimension of 1.2 m, the probability that a lig ht photon will pass through more than 16 cells is extremely low. For a cell density of 3.2106 cells/L the measured s was null. Proceeding as above we obtain that in this case the average distance of the

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160 cells (considered as points) is dc = 1mm / 3 c = 6.8 m which th erefore indicates a close-packed structure in the real case and a practical impossibility for photons to reach the semiconductor surface. Figure 5.12. Schematic representation of th e scattering and absorption of photons by the hemoglobin contained in the cell. Summarizing, the cell concentration dependent s measured in § 5.4 is, unfortunately, not caused by the electrical charge of cells but by the light scattering and absorption operated by hemoglobin. While the effect of cells on semiconductor surfaces may be present, other methods for measuring this interface electric ally are needed and will be discussed in the next chapter. 5.7. Summary In conclusion we have reported pioneering and challenging work that has been done to this point to try to investigate the elec tronic interactions betw een a semiconductor and a biological cell. We have described the experimental procedures developed for

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161 performing CPD measurements of semiconducto r-electrolyte systems and we have shown that, thanks to an appropriate calibration of the system a nd to suitable experimental approaches, CPD measurements of semiconduc tors in liquid are not only possible but highly repeatable. Despite the success reached in the preparatory phase necessary for the implementation of CPD measurements of se miconductor-cell-electrolyte systems, the results obtained from the final measurements did not yield the desire d results. It appears that the electronic effect of cell charge on the semiconducto r energy bands is lower in magnitude than expected and therefore not de tectable with the implemented CPD system which displayed a maximum accuracy of roughl y 15 mV. In addition we have developed a theoretical model that iden tifies a major issue in the cell adhesion morphology: since cells are in strict contact with the se miconducting surface only in the proximity of adhesion plaques, the quantity of charge that in a liquid environment, can affect the electronic status of the semiconductor is dras tically reduced. For CP D investigations of RBC-SiC systems, where a stronger cell elect ronic effect might have been expected, optical issues complicated the matter making it impossible to probe the effect of cell charges on the semiconductor el ectronic state with a method utilizing illumination. A careful examination of the results and insigh tful determination of their possible causes kept us from mistakenly associating the obs erved promising results with the effect of erythrocytes’ charge on the se miconductor surface potential. We believe that the results presented in this section, revealing many of the possible problems that may be encountered when studying cell-semiconductor systems via CPD, may be extremely helpful for the future im plementation of improved apparatus suitable for cell-semiconductor electronic interaction in vestigations. In fact, we do not exclude a

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162 priori the possibility that, by increasing the measuring sensitivity of the CPD apparatus used (e.g., lowering the average standard devi ation of measurements in liquid below 15 mV)) and by using cells that disp lay a higher charge than the ones used in this work, an effect of the cell charge on the electronic status of the semiconductor substrate may be monitored. However, since still very few data presently exist regarding the surface charge of biological cells, and since the surface charge is dependent on the media where cells are suspended, it may be particularly difficult to select the right cell-media combination which may yield successful results. CPD monito ring of electrogenically active cells (e.g., neurons) cultured on SiC surfaces may be a study approach that could possibly lead to the detection of measurable electronic signals (e.g., observation of the electric extracellular ‘spikes’ which offers the possibility of an AC measurement tech nique and thus higher signal to noise ratio). Also, other techniques ma y be tried in the future to study the still unknown cell-semiconductor electronic interacti ons. However, this interesting matter goes beyond the purpose of this chapter and hence w ill be exhaustively treated in Chapter 6.

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163 Chapter 6. Conclusion and Future Work 6.1. Conclusion This work is an extensive study which inves tigates the properties of crystalline SiC (e.g., 3C-, 4H-, and 6H) for bio-sensing app lications. In Chapter 2, we have described how we prepared SiC surfaces for mor phological, chemical and crystallographic characterization by using H-etching. An optim um etching process has been developed for each SiC polytype, yielding well-ordered, atomi cally flat surfaces perfectly suitable for surface science studies. This a llowed us to characterize in depth the surfaces of all the studied polytypes and in particular H-etched 3C-SiC surfaces that have displayed, upon LEED analysis, a surface reconstruction (e.g., (51)) never investigat ed in the past. The SiC surfaces presented and char acterized in Chapter 2 are extremely appealing for bioresearch applications. In fact, several biom olecular surface science studies are presently investigating the interaction of cell proteins with well-prepared, atomically flat surfaces for achieving a better unders tanding of the semiconductor/ cell interface [139, 140]. In Chapter 3 we have shown, for the first time, that single-crystal SiC is biocompatible and capable of directly in terfacing cells without the need for surface f unctionalization. Also, SiC has been shown to be signifi cantly better than Si as a s ubstrate for cell culture, with a noticeably reduced toxic effect and enhanced cell proliferation. This result opens up exciting perspectives for the use of SiC in bio-technological applica tions, suggesting that SiC should be preferred to Si which, at pres ent, is the leading cr ystalline semiconductor

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164 for bio-applications. The main factors that have been shown to define SiC biocompatibility are its hydrophilicity and surface chemistry. SiC surface morphology is shown to influence cell adhesion only wh en macropatterned while SiC polytypism and doping concentration seem to have no influe nce on cell proliferation. We have also brought to attention how the cleaning chemistry may affect cell proliferation and emphasized the importance of the selection of an appropriate cleani ng procedure for biosubstrates. From the results reported in Chapter 3 it can be easily conc luded that SiC is an ideal substrate for bio-applications such as smart-implants, drug delivery and cellular electronic interaction studies. The latter possibi lity has been particular ly investigated in Chapters 4 and 5 leading to interesting results. Specifically, in Chapter 4, a CPD apparatus has been implemented and calibra ted for measurements of semiconductors immersed in liquid, that in the past have been shown to be particular ly challenging. This measurement apparatus has been used to charac terize the electronic status of SiC surfaces upon different chemical charging processes and has also been used to investigate the effect that H-etching has on th e electrical propertie s of SiC polytypes (e.g., measurements in air, Chapter 4). From these studies we have obtained an accurate description of the response of SiC surfaces to added charges and the novel result that H-etching electronically passivates 3C-SiC(001) surface s. The latter finding may be a topic for future studies both in the surface science and in the bio-medical fields. In Chapter 5, CPD measurements of SiC surfaces immersed in electrolytes have been performed and the obtained results have been discussed with the ai d of existent theoretical models to define the electronic effect of different elect rolytes on SiC surfaces. Unfortunately, CPD measurements of semiconductor-cell-electroly te systems did not reveal a measurable

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165 effect of mammalian cell charge on the electr onic status of SiC surfaces. However, this acquired knowledge has allowed us to creat e a model where cell adhesion morphology and electric field decay are of primary im portance in defining the amount of charge actively influencing the semiconductor surface. B ecause of this model, it appears that the cell charge effect on semiconducting surfaces is smaller in magnitude than initially estimated and cannot be detected with a m easuring system which presents a maximum accuracy of 15 mV. Summarizing, this work exhaustively i nvestigated SiC’s surf ace characteristics and bio-potentialities and for the first time introdu ced crystalline SiC as a biomaterial. At the same time a pioneering first step which help s to move towards a better understanding of cell-semiconductor electronic interactions was made, leading to interesting results that may help to select new approaches for su ccessful future investigations (§ 6.2). 6.2. Future work The impossibility of detecting the charge associated with cells while performing CPD measurements (§ 5.4-6) has been e xplained by a suitable model (§ 5.6) that identifies cell adhesion morphology and cell char ge electrical decay as two of the major causes. However, the initial idea of detect ing, using a contactle ss technique, the band bending induced by cells in semiconducting surfaces is still extremely appealing and should not be abandoned. Possible modifications that could be made to the presented CPD measurements and hence lead to successful results are as follows: 1) elimination of the deep illumination feature that has b een shown to complicate the measurement (Chapter 5); 2) monitoring of adherent electr ogenically active cells such as neurons, heart muscle cells, or toad bladder epithelial cells [122, 141, 142] since their elect rical activity

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166 has been well characterized by previous studi es [132]. Hence, time-dependent monitoring of the electrical activity of these cells and its influence on the semiconductor electronic state could be attempted. The observation of the effect of an AC signal on the semiconductor space charge would be most likel y facilitated by the hi gher signal to noise ratio of the experiment. Also, completely different approaches may be attempted to investigate the electronic effect of cells on semiconducti ng surfaces. One possibility would be, for example, to use XPS measurements of semiconducting surfaces whose molecules are bonded to adhesion proteins to define the entity of band bendi ng introduced by the biological matter. In fact, as already discussed in Chapter 4, core level binding energies in XPS spectra vary with the energy band bending and can be us ed to quantify energy band shifts. Furthermore, XPS studies coul d be performed to identify the chemical components of the adhesion proteins remaini ng on sample surfaces after trypsinization or RCA cleaning (Chapter 3). We actually have already started, in collaboration with Dr. Starke’s group (Max Planck Institute, St uttgart, DE), the aforementioned XPS investigations which have prov ided promising preliminary re sults. Specifically, two 3CSiC substrates were used in these prelimin ary experiments. HaCa T human keratinocytes cells were cultured on one, wh ile the other was used as control. Both samples were immersed in media during cell culture and th en the cells were trypsinized (Trypsin + EDTA, (ATCC)) after a thorough rinse in PBS. Subsequently, the samples were immersed in DMEM culturing mdia (10 % of FBS added) to block the trypsinization process and leave part of the a dhesion proteins attached to th e substrate, and then rinsed in DI water. XPS analysis of the two surf aces provided interesting results. The sample

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167 with adhesion proteins displayed a hi gher content of carbon with several unknown components (see Fig. 6.1). Figure 6.1. XPS spectra of 3C-SiC sa mples with and without cell adhesion proteins.Several unknown C-peaks are present in both the spectra. No comparison should b e made between the peak magnitudes of the two samples; only elemental percents within an individual sample are significant. The identification of the different car bon chemical components present on the spectra may yield precious information re garding the binding m echanism between SiC and mammalian cells. Also, in-depth studies of core binding energy levels for these samples may lead to the quantif ication of energy band shifts induced by the presence of biological material. The interesting passivating effect that hydr ogen etching has on 3C-SiC surfaces will surely be a rich topic for future investigations. Scanning tunneling microscopy (STM) could be used to finalize a model to descri be the surface structure of H-etched 3CSiC(001) samples. Also, additional ATR-FTIR or infrared abso rption spectroscopy (IRAS) studies may be performe d to investigate in depth the molecular binding present at

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168 this surface which are a probable cause of the observed electr onic passivation [143]. Since it has been suggested that transfer of electrons from the occupied valence band of proteins into free stat es of the semiconductor may cause protein modifications [30, 12], the electronic passivation observed on etched 3C-SiC surfaces may lead to different enlightening experiments. Prot eins (e.g., fibrinogen) could be deposited on electronically passivated and standard 3C-SiC surfaces. A di rect comparison of the protein behavior (e.g., if electronic exchange occurs, fibri nogen decomposes into fibrin monomer and fibrino-peptide) would allow de tection of differences in the electronic behavior of the two substrates. If, for example, proteins do not decompose on the H-etched surfaces (as might be expected due to the electronic passivati on), H-etched 3C-SiC surfaces would represent an ideal material for blood sensing devices. It is in fact know n that a high charge exchange between blood proteins and contac ting surfaces leads to thrombo-formation. Zeta potential measurements of SiC surf aces could also be attempted to better understand the electronic behavi or of this material within electrolytes. The results obtained from these studies w ould be of fundamental importa nce for a better definition of the potentiality of this material for in-vivo applications and may facilitate future cellelectronic interactions studies.

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About the Author Camilla Coletti was born on November 9th 1979 in Marsciano (Perugia), Italy. In 1998, she earned the Maturita’ Scientifica from the Scientific Liceum ‘L. Salvatorelli’ of Marsciano (IT). In 2004 she received the Laurea (M.S. degree) cum laude in electrical engineering from the University of Perugia (IT) with a thesis entitled “Characterization of Scanning Spreading Resistance Microsco py (SSRM) measurements by device simulation” performed in collaboration with the Swiss Federal Institu te of Technology (ETH) of Zurich (CH). In 2005, she passed the Esame di Stato (italian Professional Engineer qualification) and entered the el ectrical engineering Ph.D. program at the University of South Florida (USF) in Tamp a, FL (USA). During her Ph.D. program she has collaborated with the Max-Planck Institute (MPI) of Stuttgart (DE) and the Institute of Photonics and Nanotechnology of the Italia n National Research Council (IFN-CNR) of Trento (IT) and authored two publications in Applied Physics Letters (one first author) and several peer-reviewed conference papers.