SYNTHESIS and NANOENGINEERING of GALLIUM PHOSPHIDE NANOSTRUCTURES for PHOTOELECTROCHEMICAL SOLAR ENERGY CONVERSION by Wen Wen A

SYNTHESIS AND NANOENGINEERING OF GALLIUM PHOSPHIDE NANOSTRUCTURES FOR PHOTOELECTROCHEMICAL SOLAR ENERGY CONVERSION

Wen Wen

A dissertation submitted in partial fulfillment of the requirements for the degree of Doctor of Philosophy (Chemistry) in The University of Michigan 2014

Doctoral Committee:

Assistant Professor Stephen Maldonado, Chair Assistant Professor Bart Bartlett Associate Professor Peicheng Ku Professor Adam Matzger

DEDICATION

For my parents, my husband and my daughter.

ACKNOWLEDGMENT

First, I would like to thank my advisor, Professor Stephen Maldonado. Thank you very much for your patient guidance during my last five years. I want to thank you for taking time for one-to-one weekly meeting with me for 5 years and preparing me for each important stage in graduate school (seminar, candidacy, data meeting and defense). I also would like to thank my committee member, Professor Adam Matzger, Professor Bart Bartlett and Professor Peicheng Ku. Dr. Matzger, I started my graduate school by rotating in your group. It is an invaluable experience. And I also want to thank you so much for letting using instruments in your lab. Dr. Bartlett, I took Inorganic Chemistry with you, I learned a lot from your class and it is very helpful in my PhD. research, thank you. Dr. Ku, thank you very much for your feedback and advice for my research during our meeting. Looking back to my past five years in graduate school, I think I could not feel any luckier to know my labmates. Thank you very much for your help and support. Sean Collin, for taking me to numerous SEM and TEM sections and you are always very patient to answer my question. Azhar Carim, for helping me with characterization of my samples and showing me a lot of tricks. Dr. Michelle Chitamber, for helping me out with the spectral response, even you left, your promote reply for my questions is very helpful. Dr. Jeremy Feldblyum, for your friendship and you are always there to encourage me. Sabrina Peczonczyk and Junsi Gu, my peers for five years, we went through each milestone of Ph.D life together and thank you so much for your support. Eli Fahrenkrug, for being a nice cubemate and your helpful advices in research and life. Besty Brown, for always being very sweet, no matter your candies or your kind comforts. Luyao Ma, Sudarat Lee and Tim Zhang, for taking lab duties and your support. This dissertation could not be done with the help from the department staffs. I would greatly thank Roy Wentz from glass shop, you are always there whenever I need to try new setups and offer me countless good ideas to make them happen. Stephen

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Donajkowski from machine shop, for always helping me fixing instruments promptly. Dr. Haiping Sun and Dr. Kai Sun from EMAL, for your generous help in solving technical problems. I would like to take this opportunity to thank my previous advisors, Professor Ranko Richert from Arizona State and Professor Kai Wu from Peking University, for your patient guidance and great support. Last but not least, I would like to thank my family. My parents, for giving me a family surrounded with love and raising me up to who I am with endless love and support. My husband Chen Ling, the love of my life. You are always there supporting me through all the frustrations and challenges I met. I could not imagine I can get to the point of what I am now without you. I am looking forward to share the journal of the rest of my life with you.

TABLE OF CONTENTS

DEDICATION...... ii ACKNOWLEDGMENT ...... iii LIST OF FIGURES ...... viii LIST OF TABLES ...... xiv ABSTRACT ...... xv CHAPTER I Introduction ...... 1 I.1 Importance of Solar Energy ...... 1 I.2 Solar Energy Conversion ...... 1 I.2.1 Photosynthetic Energy Storage ...... 2 I.2.2 Photovoltaic Energy Conversion ...... 2 I.2.3 Photoelectrochemical (PEC) System ...... 6 I.3 Gallium Phosphide Nanowires for PEC Application ...... 19 I.4 Content Description ...... 20 I.5 References ...... 23 CHAPTER II Photoelectrochemical Behavior of n-type Gallium Phosphide Nanowires as Photoanodes ...... 28 II.1 Introduction ...... 28 II.2 Methods ...... 29 II.2.1 Preparation of GaP Nanowires ...... 29 II.2.2 SEM/TEM Characterization of GaP Nanowires ...... 29 II.2.3 Diffuse Reflectance Measurement ...... 29 II.2.4 Photoelectrochemistry ...... 32 II.3 Results ...... 33 II.3.1 Morphology of GaP Nanowries ...... 33 II.3.2 Photoelectrochemical Response of as-Prepared Nanowire Films ...... 33

II.3.3 Diffusive Reflectance Measurement of GaP Nanowire Films ...... 38 II.3.4 Spectral Response of Nanowires in Different Electrolytes ...... 38 II.3.5 Effect of Catalyst Residue in the Overall Quantum Yield Performance ...... 40 II.4 Discussion ...... 42 II.5 Conclusions ...... 46 II.6 References ...... 46 CHAPTER III Structural and Photoelectrochemical Properties of Gallium

Phosphide Nanowires Annealed in NH3 ...... 49 III.1 Introduction ...... 49 III.2 Methods ...... 51 III.2.1 GaP Nanowire Film Preparation ...... 51 III.2.2 Materials Characterization ...... 51 III.2.3 Photoelectrochemical Measurements ...... 55 III.3 Results ...... 55 III.3.1 X-ray Diffraction ...... 55 III.3.2 Diffuse Reflectance ...... 57 III.3.3 Raman Spectroscopy ...... 59 III.3.4 Transmission Electron Microscopy ...... 65 III.3.5 X-ray Photoelectron Spectroscopy ...... 69 III.3.6 Long Wavelength Photoelectrochemical Conversion ...... 72 III.4 Discussion ...... 74 III.5 Conclusions ...... 77 III.6 References ...... 78 CHAPTER IV Zinc Doping of Gallium Phosphide Nanowire Films ...... 82 IV.1 Introduction ...... 82 IV.2 Methods ...... 83 IV.2.1 Zinc Doping Methods ...... 83 IV.2.2 Photoelectrochemical Measurements ...... 89 IV.3 Results ...... 93 IV.3.1 Zinc Doping ...... 93 IV.5 Conclusions ...... 109

IV.6 References ...... 110 CHAPTER V Conclusions and Future Work ...... 111 V.1 General Findings ...... 111 V.2 Future Directions ...... 112 V.2.1 P-type Doping of GaP Nanowires ...... 112 V.2.3 GaP Nanowires for PEC Cell ...... 117 V.3 References ...... 118

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LIST OF FIGURES

Figure I.1 A schematic overview of photosynthesis process. (Source: http://en.wikipedia.org/wiki/Photosynthesis) ...... 3 Figure I.2 (a) A schematic band structure of an intrinsic semiconductor with Fermi-level aligned in the center of the bandgap;(b) A schematic band structure of an n-type semiconductor with Fermi-level aligned close to conduction bandedge;(c) A schematic band structure of a p-type semiconductor with Fermi-level aligned close to the valance bandedge...... 5 Figure I.3 (a) The charge distribution at the interface of p-n junction; (b) The band diagram of a p-n junction...... 7 Figure I.4 A schematic graph of three solar energy conversion methods. The left box represents the photosynthesis, the middle box represents the photoelectrochemical cell and the right box illustrates the photovoltaic cell...... 8 Figure I.5 Schematic depictions of a) flat band, b) accumulation, c) depletion, and d) inversion for an n-type semiconductor/liquid junction under equilibrium. (top) Energy band diagrams of the semiconductor/liquid contact. (bottom) Charge distribution in the bulk semiconductor and at the semiconductor/liquid interface...... 10 Figure I.6 Illustrations of two types of PEC cell: (a) Regenerate PEC cell with p-type semiconductor as photocathode; (b) Photosynthetic cell in the case of water splitting.... 12

Figure I.7 A schematic graph of a I-V curve with four quantities described, Isc (short circuit current), Voc (open circuit voltage), Pmax (maximum power point) and ff (fill factor), here the blue rectangle represents the product of maximum power point voltage and current...... 14 Figure I.8 Schematic depiction of primary bulk and surface recombination processes for an inorganic semiconductor photoelectrode: (1) Radiative recombination; (2) SRH recombination; (3) Auger recombination; (4) Surface SRH recombination; (5) Heterogenous charge carrier recombination...... 16 Figure I.9 Band positions of several inorganic semiconductors in contact with aqueous electrolyte at pH=1. The lower edge of conduction band and upper edge of valence band are represented by red and green respectively.(Source: Michael Grätzel, Nature,2001,414,338-344) ...... 18

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Figure I.10 Schematic depiction: (a) Carrier collection on the semiconductor/electrolyte junction for bulk semiconductor; (b) collection on the semiconductor/electrolyte junction for semiconductor nanowires...... 21 Figure II.1 a) A schematic graph of the chemical vapor deposition system used for direct sublimation of source material; b) A schematic graph of Vapor-Liquid-Solid (VLS) nanowire growth, including three steps: I) Alloying; II) Nucleation; III) Axial growth. 30 Figure II.2 Photoelectrochemical cell (PEC) set up with a flat bottom quartz cell and three electrode system. The tested sample is made as a working electrode.A Platinum coil is used as a counter electrode. A Luggin capillary with platinum inside is the reference electrode...... 31 Figure II.3 a) Cross-section SEM image. scale bar: 30 μ m ; b) Diameter distribution of as-prepared GaP nanowire film; c) Low magnification transmission electron microscope image of GaP nanowires, scale bar: 100 nm; Red arrows indicate the positions of twining defects; d) High resolution TEM image of single GaP nanowire, scale bar: 5 nm inset: selected area electron diffraction pattern with [110] beam direction...... 34 Figure II.4 J-E behavior of GaP nanowire film in dry acetonitrile with 0.5 mM ferrocenium, 49.5 mM ferrocene and 1M LiClO4. The black curve represents the J-E response in the dark. The red curve represents the J-E characterization under 100 mW·cm-2 of simulated AM 1.5 white light illumination, inset: zoomed in J-E curve in the dark...... 35 Figure II.5 Quantum yield measurement as a function of wavelength of GaP nanowires in dry acetonitrile with 0.5 mM Ferrocenium, 49.5 mM Ferrocene and 1 M LiClO4. The black line represents the photoresponse of p-type GaP crystalline wafer. The other five curves represent five photoelectrodes of GaP nanowire films prepared under the same condition (T=800 oC, t =30 min) in current CVD setup...... 36 Figure II.6 A plot of diffusive reflectance as a function of wavelength of as-prepared GaP nanowire film grown on p-GaP wafer...... 37 Figure II.7 Quantum yield measurement of GaP nanowire photoelectrodes in two different electrolytes: Ferrocene (49.5 mM) / Ferrocenium (0.5 mM) with 1M LiClO4 in acetonitrile nonqueous solution and 1 M KCl aqueous solution. The black curve represents the photoresponse of GaP nanowire photoelectrode in light yellow nonaqueous electrolyte, the red curve represents the photoresponse of GaP nanowire photoelectrode in transparent electrolyte...... 39 Figure II.8 Quantum yield measurement of GaP nanowire photoelectrodes treated with KI/I2 solution for different time duration. The black dashed line represents the photoresponse of as-prepared GaP nanowire photoelectrode. The red line represents the photoresponse of GaP nanowire photoelectrode treated with KI/I2 for 1 min. The blue line represents the photoresponse of GaP nanowire photoelectrode treated with KI/I2 for 15 min...... 41

Figure II.9 A comparison of GaP nanowires prepared by our CVD method and reported GaP nanowires prepared by solution synthesis. The black line represents the photoresponse of our synthesized sample and the error bars illustrates the deviation of five samples. The red line represents the quantum yield of the reported GaP nanowire. (The red curve is reproduced from the work: Liu, C. et.al. “Zn-Doped p-Type Gallium Phosphide Nanowire Photocathodes from a Surfactant-Free Solution Synthesis”. Nano Lett. 2012, 12 (10), 5407-5411.) ...... 43

Figure III.1 Optical micrographs depicting (a) the GaP nanowire and (b) GaP1-xNx nanowire from which the polarized Raman spectra presented in Figures III.4b and III.4c were collected, respectively.(The notation describing the polarization conditions employed in figures 3b and 3c follow the convention employed by Pauzauskie et al.19 wherein the terms to the left and right of the parentheses denote the propagation directions of the incident and collected beams, respectively, and the first and second terms in parentheses denote the polarization directions of the incident and collected beams, respectively. For example, the term indicates that the excitation is incident in the Y direction and is polarized along the X axis and the Raman scatter is collected with polarization along the Z axis in the opposite Y direction.) ...... 54 Figure III.2 (a) Cross sectional view of a scanning electron micrograph of a GaP nanowire film grown on Si (100) substrate. (b-d) X-ray diffractograms of GaP nanowire films annealed in (b) pure NH3(g) at T = 800 °C, (c) pure NH3(g) at T = 750 °C, and (d) NH3(g):Ar(g) (1:5) at T = 800 °C...... 56 Figure III.3 (a) Plan view optical images of GaP nanowire films as shown in Figure III.2a after annealing treatments at T = 750 °C with various concentrations of NH3(g). (b) Spectral profiles of the diffuse reflectance of light collected from the top of dry GaP nanowire films in air. The designations denote the annealing conditions for each sample, as in Figure III.3a...... 58 Figure III.4 (a) Representative Raman spectra for GaP nanowire films collected with excitation at 514.5 nm. The spectra are offset for clarity and normalized to the intensity of the phonon mode at 365 cm-1. The insets correspond to the same spectral data magnified by a factor of 20. Spectra were collected without any excitation or collection polarizing optics. (b) Polarized Raman spectra of an individual GaP nanowire cast on a glass slide. Light was incident and collected at the surface normal (Y vector). The terms in the parentheses indicate the polarization direction of the incident and collected light, respectively. For these measurements, the long axis of the nanowire was parallel to the Z- axis. The spectra are offset vertically for clarity and have been rescaled by the designated factor for ease of comparison. (c) Same as in (b) except for a GaP nanowire after annealing in a gas stream of NH3(g) and Ar(g) (1:5, v/v)...... 60 Figure III.5 (a)-(d) show Raman spectra of representative gallium phosphide nanowire films highlighted in different wavenumber ranges...... 62 Figure III.6 Raman spectra of GaP nanowires film annealed at 750 °C in a gas stream of NH3 (g) and Ar (g) (1:5, v/v) with excitation at different laser wavelength. The red line

x describes the Raman spectrum obtained with 632.8 nm HeNe laser. The black line shows the Raman spectrum with 514.5 nm Ar+ laser as excitation source...... 64 Figure III.7 Transmission electron micrographs of representative GaP nanowires after annealing at T = 750 ºC. (a, f) GaP nanowire without annealing; (b,g) NH3(g):Ar(g)(1:20 v/v); (c,h) NH3(g):Ar(g) (1:14 v/v); (d,i) NH3(g):Ar(g) (1:9 v/v); (e, j) NH3(g):Ar (g) (1:5 v/v)...... 66 Figure III.8 Transmission electron micrographs of a representative GaP nanowire annealed in pure argon at 750 °C for 4h: (a) low magnification and (b) high resolution micrographs, inset: selected area electron diffraction pattern...... 67 Figure III.9 High resolution transmission electron micrograph of a representative GaP nanowire annealed in a gas stream of NH3 (g) and Ar (g) (1:1, v/v), two pararell lines with arrows describe local lattice orientation...... 68

Figure III.10 High resolution N1s X-ray photoelectron spectra obtained with GaP nanowire films annealed at T = 750°C (bottom) in pure Ar(g), (lower middle) in a gas stream of NH3(g) and Ar(g) (1:9, v/v), and (upper middle) in a gas stream of NH3(g) and Ar(g) (1:1, v/v). (Top) High resolution spectra from a GaN thin film. Vertical dotted lines indicate the binding energy peak positions for the top three spectra. Spectra are offset vertically for clarity...... 70

Figure III.11 High resolution Ga3d X-ray photoelectron spectra obtained with GaP (111)B planar wafer (bottom), GaP nanowire films without annealing (middle) and GaP nanowire films annealed at T = 750°C in a gas stream of NH3(g) and Ar(g) (1:5, v/v) (top)...... 71 Figure III.12 (a) Comparison of the wavelength dependence of the external quantum yields obtained with (open diamonds) a planar n-GaP and (closed circles) a n-GaP nanowire film photoelectrode immersed in a 1M LiClO4 acetonitrile solution containing 5 mM ferrocene (Fc) and trace amounts of ferrocenium (Fc+). The data were recorded at short circuit (i.e. E = 0 V vs E(Fc+/Fc)). (b) Wavelength dependence of the external quantum yields obtained with GaP nanowire films after annealing at T = 750°C with dilute NH3(g). This spectral range corresponds to wavelengths below the lowest bandgap energy of pure GaP. Inset: the external quantum yield at λ=600 nm plotted as a function of the concentration of NH3(g) used for the annealing step...... 73 Figure IV.1 A schematic graph of the open tube system. In the top part of the scheme, the mass flow controller used to control the flow of incoming carrier gas is connected to one end of the tube. The other end of the tube is connected to the outlet of the gas. In the bottom part of the graph, the details of the placement of the zinc source and the GaP nanowire film are described...... 84 Figure IV.2 A schematic graph of spin-on dopant procedure, including four steps: (1) coat wafer with spin-on glass; (2) bake the wafer to drive off solvents; (3) high temperature drive-in anneal; (4) remove glass layer with HF...... 85

Figure IV.3 A scheme of evaporation of zinc as source: (1) One side of GaP wafer was evaporated with a layer of zinc; (2) The evaporated GaP side facing towards GaP nanowire film was placed on top of GaP nanowire film by two spacers on two sides. .... 88

Figure IV.4 (a) SEM image of GaP nanowires with the source mixture of Zn3P2:GaP (1:10 by mass); (b) Corresponding EDS of the nanowires of (a); (c) SEM image of GaP nanowires with source mixture of Zn3P2: GaP (62:85 by mass); (d) Corresponding EDS of the nanowires of (c)...... 90

Figure IV.5 (a) SEM image of GaP nanowires with the source mixture of ZnP2:GaP 1:9 by mass); (b) Corresponding EDS of the nanowires of (a); (c) SEM image of GaP nanowires with source mixture of Zn3P2: GaP (1:5 by mass); (d) Corresponding EDS of the nanowires of (c)...... 91 Figure IV.6 (a) J-E characterization of GaP nanowires prepared by in-situ doping on degeneratively doped p-Si wafer with the mixture of Zn3P2 and GaP powder (1:10 by mass); (b) J-E characterization of GaP nanowires prepared by in-situ doping on degeneratively doped p-Si wafer with the mixture of ZnP2 and GaP powder (1:9 by mass)...... 92 Figure IV.7 (a) J-E characterization of GaP nanowires on degenerately doped p-Si prepared by ex-situ doping in closed end tube system ( 9.6 mg ZnP2, GaP nanowire film 1inch away from source, P= 60 mTorr, T = 650 oC, t = 20 min); (b) J-E characterization of GaP nanowires on degenerately doped p-Si prepared by ex-situ doping in closed end tube system ( 9.6 mg ZnP2, GaP nanowire film 1inch away from source, P= 60 mTorr, T = 600 oC, t = 20 min)...... 95 Figure IV.8 J-E characterization of GaP wafer by ex-situ doping in open end tube system (a) 45.0 mg Zn, non-doped GaP wafer 7 inch away from source, T = 700 oC, t = 90 min, Ar flow: 150 sccm; (b) 25.0 mg ZnP2, non-doped GaP wafer 5.75 inches away from source, T = 800 oC, t = 90 min, Ar flow: 100 sccm; (c) 45.0 mg Zn, non-doped GaP wafer 5.75 inches away from source, T = 800 oC, t = 90 min, Ar flow: 100 sccm; (d) 45.0 mg Zn, non-doped GaP wafer 5.75 inches away from source, T = 800 oC, t = 120 min, Ar flow: 50 sccm...... 98 Figure IV.9 J-E characterization of GaP nanowires on p-GaP prepared by ex-situ doping in open end tube system (a) 16.2 mg ZnP2, GaP nanowire film 1inch away from source, T = 700 oC, t = 60 min, Ar flow: 20 sccm ;(b) corresponding quantum yield measurement of GaP nanowires prepared by the condition in(a); (c) 20.1 mg ZnP2, GaP nanowire film o 1inch away from source, T = 700 C, t = 90 min, Ar flow: 20 sccm; (d) 20.0 mg ZnP2, GaP nanowire film 1inch away from source, T = 750 oC, t = 60 min, Ar flow: 20 sccm. 99 Figure IV.10 GaP nanowires on p-GaP were treated with different amount and times for spin coating, then were preannealed at 140 oC for 60 min with rate at 500 sccm Ar, and heated at 700 oC for 120 min with 200 sccm Ar. a) J-E characterization of GaP nanowires with 62.5 μl/time for 16 times; b) Quantum yield measurement of the sample corresponding to a), inset: SEM image of the prepared nanowires; c) J-E characterization of GaP nanowires with 62.5 μl/time for 24 times; d) Quantum yield

xii measurement of the sample corresponding to c), inset: SEM image of the prepared nanowires; e) J-E characterization of GaP nanowires with 62.5μl/time for 32 times; f) Quantum yield measurement of the sample corresponding to e), inset: SEM image of the prepared nanowires...... 103 Figure IV.11 (a) J-E characterization of GaP nanowire film on p-GaP with a silicon wafer (1 mm away) facing down treated at 650 oC for 2 h under vacuum. The Si wafer is spin coated with spin on glass 100 μl /time for 10 times, (b) Corresponding quantum yield measurement of the sample prepared by (a); (c) J-E characterization of GaP nanowire film annealed at 650 oC for 2h under vacuum with the Si wafer (1mm away) on the top. Si substrate was evaporated with a layer of zinc (100 nm), before annealing, GaP nanowire film was etched by gold etchant for 60 min; (d) Corresponding quantum yield measurement of the sample prepared by (c); (e) J-E characterization of GaP nanowire film prepared at the same annealing condition as (c). The difference is that the GaP nanowires were not etched; (f) Corresponding quantum yield measurement of the sample prepared by (e)...... 105 Figure IV.12 (a) Optical image of a GaP nanowire film in an sealed ampoule with 0.020 g Zn powder under 1 atmosphere of N2; (b) Optical image of the GaP nanowire film of (a) annealed at 600 oC for 60 min; (c) J-E characterization of annealed GaP nanowire film; (d) Corresponding quantum yield measurement of GaP nanowire film of (c)...... 107 Figure V.1 (a) A schematic sketch of the cross section of a single nanowire field effect transistor device; (b) A schematic sketch of a single nanowire field effect transistor device from top view...... 113 Figure V.2 (a) Dye-sensitized hole injection from a photoexcited chromophore at the surface of a phosphide semiconductor. (b) Dye-sensitized hole injection from a photoexcited chromophore at the surface of a phosphide semiconductor under depletion ...... 115

Figure V.3 A schematic graph of p-type dye sensitized photocathode for H2O reduction reaction...... 116

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LIST OF TABLES

Table III.1 Raman spectral features for GaP nanowires annealed at T = 750°C...... 63

Table IV.1 Mass ratio of zinc source and GaP ...... 83 Table IV.2 Summary of Conditions for Zn Doping of a GaP Wafer in an Open Tube System...... 86 Table IV.3 Summary of Conditions for Zn Doping of GaP Nanowire Films in an Open Tube System...... 86

Table IV.4 Conditions for Zn Doping in an Ampoule...... 89

Table IV.5 Results summary of Zn Doping of a GaP Wafer in an Open Tube System. .. 97 Table IV.6 Results summary of Zn Doping of GaP Nanowire Films in an Open Tube System...... 97 Table IV.7 Results summary of different conditions of Spin on Glass procedures for GaP nanowires...... 102

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ABSTRACT

Gallium phosphide (GaP) is a technically mature material widely used for LEDs with excellent optoelectroinic properties. It is also a good photocathode candidate for H2 and CO2 reduction in photoelectrochemical (PEC) cell. A crucial challenge lies at the center of GaP development in the PEC cell is to improve its efficiency. For thick GaP material its PEC efficiency is generally limited by the low carrier collection, while for thin GaP it is limited by the insufficient light adsorption. An effective strategy to overcome this problem is to use high aspect ratio nanostructures, for which the long axial direction allows sufficient light absorption while the short axial direction improves the carrier collection. This dissertation details synthetic methods to fabricate GaP nanostructures and the photoelectrochemical properties of GaP nanostructures. Uniform GaP nanowires averaging 150 nm in diameter and 20-30 μm in length were synthesized by direct sublimation chemical vapor deposition. The structural properties of GaP nanowires were characterized by a set of techniques, including scanning electron microscopy, transmission electron microscopy, x-ray diffraction, Raman spectroscopy and x-ray photon spectroscopy. The optical and electrical properties of GaP nanowires were further tailored by doping other elements through high temperature post-treatment. In particular, the effects of nitrogen alloying and zinc doping on the photoelectrochemical properties of GaP nanowires were comprehensively analyzed and discussed. The results demonstrate that photoelectrochemical carrier collection efficiency of GaP is remarkably improved by employing nanowire structures. By nitrogen alloying, the supra-bandgap conversion efficiency of GaP nanowires is further enhanced. Finally by zinc doping the electrical properties of GaP nanowires can be effectively controlled. This work has demonstrated GaP nanowires can be viable materials used for PEC system. As compared to other work in this area, this dissertation shows much better photoresponse than the achieved ones. It offers an efficient method to fabricate GaP nanowires as well

xv as strategies to improve their PEC efficiency. Insights from our results and direction for future work are also discussed in this dissertation.

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CHAPTER I

Introduction

I.1 Importance of Solar Energy Societal demand of energy is expected to dramatically increase.1 Presently, the main energy resources are fossil fuels, which has several disadvantages: (1) Fossil fuels are non-renewable resources; (2) Over the past 20 years, nearly three-fourths of human- 2 based CO2 emissions came from the burning of fossil fuels. The impact on the environment affects living quality,3,4 which itself has become an urgent problem. There are several possible renewable and potentially ‘clean’ energy resources including hydroelectricity,5 biomass,6 geothermal,7 tidal,8 wind9 and solar.10 As of 2012, the energy obtained domestically from all renewable resources is roughly 10.6% of the total energy used and the outlook for 2014 is just 11.3%.1 Solar energy uniquely represents a nearly unlimited energy source with 120,000 TW of power incident on Earth.10,11 Capturing a small fraction of solar energy could therefore sufficiently satisfy societal energy demands. In addition solar energy has fewer geographic constraints than other renewable energy sources. Although promising, solar energy usage is limited by cost.12

I.2 Solar Energy Conversion There are many types of systems that can convert the solar resource into a useful form of energy. Solar energy can be captured as heat and then converted to electricity.13 However, in this thesis, the focus is on the conversion of solar energy directly to electricity and/or chemicals fuels. There are three methods to convert sunlight: photosynthesis, photovoltaic cells, and photoelectrochemical cells.

I.2.1 Photosynthetic Energy Storage Photosynthesis is nature’s method to utilize solar energy. In this process, the incident sunlight is absorbed and converts carbon dioxide and water into sugar and oxygen(shown in Figure I.1). The photosynthetic reactions are traditionally divided into two stages - the "light reactions," which consist of electron and proton transfer reactions and the "dark reactions," which consist of the biosynthesis of carbohydrates from CO2. The first step is the conversion of a photon to an excited electronic state of an antenna pigment molecule.14 The electronic excited state is transferred over the antenna molecules as an exciton. The exciton, composed of a pair of electron and vacancy, could go back to the ground state and release the absorbed energy in the form of heat unless another competing chemical pathway is available. A series of electron acceptors drive the electron spatially away from its vacancy14,15 to a photosynthetic reaction center where electron transfer reactions occur that ultimately produce NADPH and ATP16 and releases a dioxygen (O2) molecule. In the dark reaction, CO2 reacts with ATP and NADPH 14,16 through Calvin cycles to synthesize the energy compound C6H12O6. In the photosynthesis process, three critical processes are critical to store solar energy in the form of chemical fuel: (1) The absorption of light by pigment molecules; (2) The spatial separation of electron/hole pairs by the reaction center; (3) The useful work by the separated carriers driving redox reaction to produce chemicals.

I.2.2 Photovoltaic Energy Conversion A brief review of some relevant semiconductor concepts is presented first, followed by a short description of photovoltaic cells.

1.2.2.1 Concepts of Merit Semiconductor Bandgap In a solid, the energy difference between the top of valance band (Evb) and the bottom of conduction band (Ecb) is defined as the bandgap (Eg,

Figure I.2a). For semiconductors relevant to solar energy technologies, Eg ranges from 0.3-3.5 eV.17,18 The bandgap of a semiconductor defines its optical absorption threshold for the photoelectric effect. That is, generation of electric charge-carrier from light absorption only occurs with photons with energies ≥ Eg. When supra-bandgap light is absorbed by a semiconductor, charges move between the valence and conduction bands.

Figure I.1 A schematic overview of photosynthesis process. (Source: http://en.wikipedia.org/wiki/Photosynthesis)

Carrier Statistics in Semiconductors In semiconductors, there are two types of charge carriers: holes and electrons. The equilibrium concentration of electrons and holes (n and p, respectively) can be expressed in the following relation.

2 E g np  ni  N C N V exp(  ) (I.1) k BT

In this equation, kB is the Boltzmann constant, T is the temperature, n and p represent the equilibrium concentration of electrons and holes, ni is the intrinsic carrier concentration, Nc is the effective density of states in the conduction band, and Nv is the effective density of states in the valence band. Generally in intrinsic semiconductors, Nc 19 -3 19 and Nv have similar values (~10 cm ). According to equation I.1, at constant temperature, even ‘small’ bandgap semiconductors have low carrier concentrations. For example, pure silicon (Si; Eg = 1.12 10 -3 V; ni = 1.5×10 cm ) has less than 1 part per trillion of free charge carriers. Since the intrinsic carrier concentration is very low, impurities that add charge carriers can have a profound impact on conductivity even at the scale of 1 part per billion. Therefore, to increase the electric conductivity of semiconductors, impurity atoms can be purposefully introduced, i.e. ‘doping’. Based on which charge carrier (electron or hole) an impurity introduces, dopants can be divided to two categories: donors and acceptors. Donor atoms readily ionize by donating electron density to the bulk semiconductor while acceptor atoms introduce additional holes in the semiconductors. Since equation I.1 holds at equilibrium, increasing the concentration of one charge-carrier type necessarily means decreasing the concentration of the other charge-carrier type. Semiconductors doped with donors are n- type and semiconductors doped with acceptors are p-type. The charge-carrier type in higher abundance is called the majority carrier and the lower concentration charge-carrier type is the minority carrier. In n-type semiconductor, the majority carrier are electrons while in p-type semiconductor, the majority carrier are holes. Fermi-level The Fermi level describes the average energy of transferable electrons and holes and is denoted as EF. Formally, the Fermi level is the energy level where the probability of being occupied by an e- is 0.5. In a pure semiconductor with no dopants at 0 K, the Fermi level locates in the middle of the bandgap. Adding dopants to a

a)Conduction b) Conduction c) Conduction Band Band Band

bandgap Fermi‐level

Fermi‐level Fermi‐level

Valence Valence Valence Band Band Band

Figure I.2 (a) A schematic band structure of an intrinsic semiconductor with Fermi-level aligned in the center of the bandgap;(b) A schematic band structure of an n-type semiconductor with Fermi-level aligned close to conduction bandedge;(c) A schematic band structure of a p-type semiconductor with Fermi-level aligned close to the valance bandedge.

5 semiconductor shifts the Fermi level. For an n-type semiconductor, the majority carriers are electrons and its Fermi level moves closer to the conduction bandedge. In contrast, p- type semiconductor has higher concentration of holes and thus its Fermi level is closer to the valence bandedge (shown in Figure I.2) The Fermi level concept only has meaning at equilibrium. When there is a 2 perturbation to the equilibrium (e.g. optical excitation), np no longer equals to ni . In this situation, there is no single energy can be used to describe the average energy of electron and holes simultaneously. Quasi-Fermi level are used to describe the average energy for electrons (EF,n) and holes (EF,p) respectively (shown in Equation (I.2) and (I.3)). n E F ,n  Ecb  k BTIn (I.2) N C N E  E  k TIn V (I.3) F ,p vb B p

I.2.2.2 Photovoltaic Cell A basic photovoltaic cell consists of two different types of solids and connected with an abrupt junction.20 A typical PV cell utilizes a p-n homojunction, i.e. a p-type semiconductor in contact with an n-type semiconductor (Figure I.3). In this case, when a photon is absorbed by the semiconductor, an electron is excited from the valence band to the conduction band and a hole is created in the valence band. The generated electron and hole pair is immediately separated by the electric field at the junction and collected at separate contacts. Current can be drawn across the contacts as usable electric power. Similar to photosynthesis, the photovoltaic cell also has a free energy gradient to separate the electron/hole pairs. However, instead of using reaction centers to separate charges, the separation is achieved by using a solid junction of two different semiconductors. Another feature of photovoltaic cell is that it can convert the solar energy to electricity.

I.2.3 Photoelectrochemical (PEC) System Since Brattain and Garret21 and subsequently Gerischer22 first studied semiconductor–electrolyte heterojunctions, photoelectrochemical (PEC) cells have been identified for renewable energy23,24. A comparison of three different solar energy

a) ‐ ‐ + + p‐type ‐ ‐ + + n‐type ‐ ‐ + +

Built‐in Electric Field b)

Conduction Band n‐type Energy EF Potential

p‐type Valance Band

Figure I.3 (a) The charge distribution at the interface of p-n junction; (b) The band diagram of a p-n junction.

Electricity Fuel

Fuel Electricity

CO2 e H2 O2

Sugar S C M H2O

H2O O2

Photosynthesis Photoelectrochemical Photovoltaic Cell Cell

Figure I.4 A schematic graph of three solar energy conversion methods. The left box represents the photosynthesis, the middle box represents the photoelectrochemical cell and the right box illustrates the photovoltaic cell.

8 conversion methods is shown in Figure I.4. In analogy to photosynthesis process, the semiconductor functions as the pigment molecules for light absorption. The electrical field near the semiconductor/electrolyte interface promotes spatial carrier separation and the separated holes or electrons are used to drive redox reaction in the electrolyte to store the energy. The free energy gradient at semiconductor/liquid junction is very similar to that in photovoltaic cell and photosynthesis. It is responsible for the charge separation and finally leads to the energy conversion. The difference is that the photovoltaic only produces electricity. I.2.3.1 Types of Semiconductor/Liquid Junctions When a semiconductor is contacted with a liquid electrolyte with a defined redox potential, heterogeneous charge transfer occurs at the interface to equilibrate the Fermi- levels throughout the interface. Figure I.5 shows four different situations (flat band, accumulation, depletion, and inversion) when an n-type doped semiconductor is inserted into a solution.25,26 Analogous scenarios can be obtained with p-type semiconductors. Each case is described below. Flat-band: When the Fermi levels of the solution and semiconductor are equivalent at the start, no heterogeneous charge transfer is necessary to reach equilibrium Accumulation: When the Fermi-level of the semiconductor is more positive than the electrochemical potential of the solution, electrons transfer from the solution to the n- type semiconductor to equilibrate the Fermi level. This results in higher electron density on the semiconductor side of semiconductor/liquid junction than in bulk semiconductor, while the residue positive charges accumulate on the interface. The spatial region where electrons accumulate is called the accumulation region and the semiconductor is under accumulation condition. A potential drop occurs inside the semiconductor between the surface and the bulk interior. The excess majority carriers are not bound to individual atoms, but instead are delocalized. The associated electric field is strong near the interface and simultaneously repels holes from the interface and draws electrons to the interface. Depletion When the Fermi level of the semiconductor is more negative than the electrochemical potential of the solution, electron transfer from the n-type semiconductor to the solution occurs to reach equilibrium. The finite number of electrons lost from the

(

a)‐ b) ( ) ‐ )

Potential Potential (+) (+)

0 0 semiconductor solution semiconductor solution

‐ ‐ ‐ + + ‐ + ‐ + ‐ + ‐ + + ‐ ‐ + + ‐ ‐ ‐ + ‐ ‐ + ‐ +

c) d) ( ‐ ( ) ‐ )

Potential Potential (+) (+)

0 0 semiconductor solution semiconductor solution

‐ ‐ + + ‐ ‐ + + + ‐ + + ‐ ‐ ‐ ‐ + + + ‐ ‐ ‐ + ‐ ‐ + + ‐

Figure I.5 Schematic depictions of a) flat band, b) accumulation, c) depletion, and d) inversion for an n-type semiconductor/liquid junction under equilibrium. (top) Energy band diagrams of the semiconductor/liquid contact. (bottom) Charge distribution in the bulk semiconductor and at the semiconductor/liquid interface.

10 semiconductor results in an increase in hole density near the surface. The depth that electrons are depleted from the semiconductor is dependent on the doping density of the semiconductor. This near-surface region is called the depletion region and the depth of this region is then called the depletion width (W). Electrical field within depletion region plays a critical role in separating carriers. Electric field effectively sweeps electrons to n-type contacts and holes to p-type contacts before they recombine in the semiconductor. Inversion A special depletion case occurs when equilibration requires enough heterogeneous electron transfer that the near surface region loses its n-type character. This is called inversion (Figure I.5d). In this case, the Fermi-level near interface can decrease below its intrinsic value, i.e. the character of interface changes to p-type while the bulk of semiconductor is still n-type. Effectively, a p-n junction forms between the near-surface and the bulk. A quantity used to characterize the semiconductor/liquid junction is the barrier height, denoted as Φb. The barrier height energy (qΦb) is defined as the energy difference between the Fermi-level of the semiconductor and the bottom of the conduction band at the interface. Based on the definition, qΦb can be expressed below:

qb  EF  Ecb (at interface) (I.4) Hence, in principle, the barrier height (and correspondingly the flat band, accumulation, depletion, or inversion condition) can be predicted once the electrochemical potential of solution and the bandedge potentials of the semiconductor are known.

I.2.3.2 Types of PEC Cells The two types of PEC cells are diagrammed in Figure I.6.27 The first type is the regenerative cell, in which the solar energy is converted to electricity through electrochemical reactions but with no net chemical change. When supra-bandgap light illuminates a p-type semiconductor electrode, the electron-hole pairs are separated by the electrical field at the interface. The majority carrier (holes) are excluded from the surface and go towards the counter electrode to oxidize the redox species from R to O at a defined electrochemical potential. At the meanwhile, the minority carrier (electrons) are

e e

a) b) e H2O

O O2 SC R M SC M H2 H O e 2 soln H O 2 Figure I.6 Illustrations of two types of PEC cell: (a) Regenerate PEC cell with p-type semiconductor as photocathode; (b) Photosynthetic cell in the case of water splitting.

12 drawn by the electrical field towards the interface and then captured by the solution to reduce O to R. When the half reactions at the semiconductor electrode and the counter electrode are the same (just different direction), no net chemical change happens in the electrolyte. Although here shows an example of a p-type semiconductor, much work of degenerative PEC cells has been focused on n-type II-VI and III-V semiconductors using 28 -29,30 both inner-sphere (e.g. sulfide/polysulfide, I2/I ) and outer-sphere (e.g. ferrocence/ferrocenium) redox couples.31 The second type of PEC cell is the photosynthetic cell and can store solar energy through the production of chemical fuels. The operating mechanism of photosynthetic cell is similar to the regenerative cell except that the photosynthetic cell involves two different redox couples at the respective electrodes. Figure I.6b shows an example where the half reactions comprise water oxidation and proton reduction, i.e. a ‘water splitting’ cell. In this cell, the action under illumination produces two new chemicals, H2 and O2, in solution. Since H2 can be reacted to release thermal or electrical energy, it represents the storage medium for the incident optical energy. Correspondingly, O2 is a not a fuel but a clean by-product that follows when water is the feedstock for H2.

I.2.3.3 Quantitative Description of the Performance of PEC Cell In general, there are three quantities that characterize the performance of PEC cell: 27 the short circuit density (Jsc), the photovoltage (Voc), and the fill factor (ff). Jsc is defined as the current that flows when a direct short circuit is present between the two electrodes under a steady-state illumination. Voc is measured at open-circuit (i..e no net current passes) under a steady-state illumination. Voc describes the maximum Gibbs free energy that can be extracted from the PEC device under a particular set of illumination conditions. Voc reflects the extent of quasi-Fermi level splitting. The fill factor is defined as the ratio of the actual maximum obtainable power to the product of the open circuit voltage and short circuit current. ff characterized the rate of current density approaches

Jsc. Graphically, it describes the “squareness” of a J-V curve. (Figure I.7)

I.2.3.4 Efficiency of a semiconductor photoelectrode The net energy conversion efficiency of a working semiconductor photoelectrode is a product of three separate aspects,32

ff = Isc ×Voc Pmax

Isc

Photocurrent 0 Voc 0 Voltage

PEC optical separation echem (I.5) where ηoptical, ηseparation, and ηechem are the respective descriptors for light absorption, charge separation, and efficiency for driving the relevant electrochemical reaction. The first term, ηoptical, describes the capacity of the semiconductor to generate an electron-hole pair with a thermalized energy equal to the bandgap energy for light absorption with energy larger than the bandgap. The second term, ηseparation, describes the internal conversion efficiency of the semiconductor for supplying photogenerated electrons and holes at the maximum possible electrochemical potential difference. This term accounts for all loss mechanisms within the semiconductor material. The third term, ηechem, relates to the electrocatalytic the electrochemical reactions with the photogenerated charge carriers. For complete analysis of the PEC efficiency, all three separate efficiency must be rigorously assessed. The optical properties of nanostructured semiconductors in the aspect of light absorption efficiency33,34,35,36 and the electrocatalysis of fuel-forming redox transformations37,38,39 have been studied extensively and are outside the scope of this thesis. This thesis focuses on understanding factors that affect the ηseparation. of nanostructured photoelectrodes. The internal quantum yield, Ф, for conversion of absorbed photons to electrons and holes collected at the contacts at steady-state conditions is also a relevant metric. For clarity, Ф is on a scale of 0 - 1. The major cause of deviation of Ф from unity comes from either the total carrier recombination in the interior volume of the semiconductor (ФR), the total carrier recombination at the interface where charge-carriers are collected 40 (ФContact), or the total carrier recombination at the surface between contacts (ФS,).

 1R Contact S (I.6) I.2.3.5 Recombination of Charge Carriers The five primary recombination processes are shown in Figure I.8. Radiative Recombination: Process 1 in Figure I.8 represents radiative recombination, which occurs by recombining a conduction band electron and a valence band hole through the emission of a photon with bandgap energy. Radiative recombination is inherent to every semiconductor device under illumination or applied bias.41 It cannot be

12345 RA‐ Conduction Band ‐ ‐ ‐ ‐ ‐ OA hν

RA‐ + + + + Valence Band OA

Figure I.8 Schematic depiction of primary bulk and surface recombination processes for an inorganic semiconductor photoelectrode: (1) Radiative recombination; (2) SRH recombination; (3) Auger recombination; (4) Surface SRH recombination; (5) Heterogenous charge carrier recombination.

16 eliminated by purification. For this reason, radiative recombination represents a fundamental limitation to the performance of an optimized semiconductor device and forms the basis for Shockley and Queisser’s limit on the maximum efficiency for optimized single-junction devices under one-sun illumination.42 Shockley-Read-Hall (SRH) recombination: Process 2 in Figure I.8 occurs when carriers recombine at an energy level inside the bandgap in a semiconductor.43,44 This trap state can originate from either defects or impurities in the semiconductor. Unlike radiative recombination, SRH recombination can be greatly decreased by purifying the semiconductor. For a SRH recombination at a single trap state, the highest level of recombination will occur when the trap energy is exactly at the middle of the bandgap. For a distribution of multiple populations of trap states with different energies, the individual recombination rate must be integrated over the entire band gap energy to determine the total recombination rate. Auger Recombination: Process 3 in Figure I.8 shows one type of three body recombination, in which two charge carriers recombine and transfer energy to the third carrier. This energy is then partially or wholly dissipated via thermalization.45 Since Auger recombination involves three bodies, this process is highly dependent on the concentration of carriers. At low carrier concentration, Auger recombination can be negligible, while Auger recombination becomes significant under high concentration of carriers such as high-level illumination. SRH recombination at Surface: Process 4 in Figure I.8 illustrates the recombination of electron and hole through a mid-gap trap on the surface of the semiconductor.32 High quality interfaces with minimal surface recombination nominally have recombination events that occur at rates less than 100 cm s-1 while heavily defective interfaces have recombination velocities in excess of 105 cm s-1. 40 Heterogeneous charge-carrier recombination: Process 5 in Figure I.8 describes charge-carrier recombination through heterogeneous charge transfer with dissolved species in solution. When both electrons and holes are transferred at the same interface, zero energy is stored in solution since there is no net chemical reaction.46,47

Figure I.9 Band positions of several inorganic semiconductors in contact with aqueous electrolyte at pH=1. The lower edge of conduction band and upper edge of valence band are represented by red and green respectively.(Source: Michael Grätzel, Nature,2001,414,338-344)

I.3 Gallium Phosphide Nanowires for PEC Application The choice of semiconductor for a PEC system is critical. The bangap of the materials highly influences the maximum photovoltage it can offer while the band positions of a semiconductor greatly affect the usability of this material for certain redox reactions. Figure I.9 describes the band position of several semiconductors in contact with different electrolytes.48

Since Fujishima and Honda first reported using n-type TiO2 for water splitting reaction,49 considerable efforts has been focused on developing suitable semiconductor in photoelectrochemical (PEC) cell to convert sunlight into chemical fuel. Numerous 50,51,52 53,54,55 56,57,58,59 materials, especially semiconductors such as Si, ZnO, TiO2, have been extensively studied. The major challenge that limits the application of these materials in practical PEC is to balance the photovoltage and the adsorption of solar irradiation. Semiconductors with wide band gaps can produce large photovoltage, but their solar adsorption efficiency is low. Similarly, semiconductors with narrow band gaps can effectively adsorb the solar irradiation, but their photovoltages are limited by the band gap. 60 Gallium phosphide has a mid-sized bandgap (Eg= 2.26 eV) that intrinsically allows for large photovoltages under illumination and still can absorb a significant fraction of solar irradiance. For instance, in the solar water splitting reaction, the thermodynamic requirement for the reaction is 1.23 V at a standard condition.49 The mechanistically complex processes for bond formation/scission in water splitting demand an additional overpotential that increases the total needed photo-induced voltage to ≥ 1.7 V.61,62 GaP has large carrier mobility (300 cm2 V-1 s-1 for electrons and 500 cm2 V-1 s-1 for holes).60 Gallium phosphide (GaP) has long been considered as a candidate semiconductor for photosynthetic photoelectrochemical applications. The band positions of GaP align well with several important redox reactions, such as H2 reduction and CO2 reduction.48 In fact, GaP was reported to be a potentially excellent photocathode material 63,64,65 in PEC system for efficient CO2 and H2 reduction in late 1970s. Since then, researchers have been working on using this material for PEC, however, several material challenges have hindered GaP-based PEC cell development.

The major drawback of using GaP in PEC system is that GaP has an indirect bandgap, which means it needs phonon activation in addition to the corresponding photon excitation.60 This leads to relatively large penetration depth (α-1) and small minority carrier diffusion length (Ld) and thus pronounced lower ηseparation. One approach of solving the problem is adapting high aspect ratio architecture to decouple light absorption and carrier collection direction (Figure I.10). In 1995, Erné et.al. reported greatly enhanced sub-bandgap photocurrent by using porous GaP.66 In 2009, Price et.al. have achieved a high performance macroporous GaP photoanode prepared by top etching GaP crystalline wafer67. However, the top-down fabrication is very costly from the economic point of view. In order to make affordable and efficient photoelectrode for PEC cell, it is essential to develop a simple and low cost preparation method. This is the main underlying motivation for the remainder of this thesis. The motivation for this dissertation is to make GaP nanowires viable photoelectrode materials for PEC cell. Nanowire features a high aspect ratio (> 20), allowing light absorption in the axial direction and carrier collection through the radial direction. The very beginning work of synthesizing nanowrie structures dated back to 1960s, when Wagner reported the growth of Si whiskers by vapor-liquid-solid mechanism.68 Since then, various other ways to grow semiconductor nanowires from vapor deposition,69,70 solution synthesis71,72 and even electrochemical deposition73 have been reported. However, to date, GaP nanowires have been predominantly synthesized by chemical vapor deposition (CVD).74,75,76,77,78,79,80

I.4 Content Description The objective of this dissertation is to study the photoelectrochemical properties of gallium phosphide (GaP) nanostructures and tune the bulk properties of GaP to make GaP nanostructures viable photoelectrode materials for solar energy conversion. Three chapters in this dissertation aim to provide a simple and low cost fabrication method to control morphology and bulk properties of photoelectrode materials. The PEC properties are studied in regenerative PEC cell. Chapter 2 describes n-type GaP nanowire photoelectrodes prepared through a simple chemical vapor deposition using Au as catalyst. The performance of these electrodes under illumination was investigated in a nonaqueous, regenerative PEC cell. Planar

a) b) Electrolyte

α -1 α -1 λ λ Semiconductor + ‐ + ‐ ‐ + Metal contact

21 electrodes featured low photon to electron efficiency in wavelength region of 450 nm to 540 nm. The minority carrier diffusion length was shorter than the optical penetration depth in GaP near the bandgap region, resulting in large bulk recombination losses. By utilizing high-aspect-ratio nanowire photoelectrodes, the processes of light absorption and carrier collection were effectively decoupled and carriers could be collected from the radial direction. This tactic resulted in a dramatic decrease in bulk recombination processes and a significant increase in the photon to electron conversion efficiency. In addition, the growth of nanowires involved using Au as catalyst and the effects of Au on the performance of nanowires are also studied.

Chapter 3 reports the utility of nanostructured GaPxN1-x in PEC energy conversion systems. Nitrogen alloyed GaP nanowires have been prepared from GaP nanowires annealed in flowing NH3-containing inert gas. A set of annealing conditions were identified to incorporate nitrogen into nanowires without formation of bulk gallium nitride. The results show the nitrogen atoms were incorporated non-uniformly throughout the nanowires by substituting the anionic sites. Increased absorption of visible light at wavelengths with less energy than the bandgap was observed in nanowires treated with higher fraction of NH3, consistent with the bandgap bowing effects in bulk GaPxN1-x alloy system. Enhancement photoresponse in the sub-bandgap wavelength region (550-600 nm) was achieved. Chapter 4 illustrates the preparation of zinc doped GaP nanowires in aqueous PEC degenerative cell. Various approaches have been attempted to efficiently dope GaP nanowires. The valuable point from doping processes is that the temperature and pressure are critical in efficient doping. Preliminary data show the GaP nanowires can be doped to be p-type by intensive treatment with Zn in a closed ampoule. Morphology is another factor to consider in the design of efficient photoelectrode. In this chapter, the growth conditions in the current chemical vapor deposition were also investigated. Chapter 5 summarizes the key findings of chapter 2, 3 and 4 and provides future work. In this chapter, the conclusions were drawn and future directions, specifically the zinc doping and morphology control are proposed.

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57. Hoang, S.; Guo, S.; Hahn, N. T.; Bard, A. J.; Mullins, C. B., Visible Light Driven Photoelectrochemical Water Oxidation on Nitrogen-Modified TiO2 Nanowires. Nano Lett. 2011, 12 (1), 26-32. 58. Zhen, C.; Liu, G.; Cheng, H. M., A film of rutile TiO2 pillars with well-developed facets on an alpha-Ti substrate as a photoelectrode for improved water splitting. Nanoscale 2012, 4 (13), 3871-3874. 59. Altomare, M.; Lee, K.; Killian, M. S.; Selli, E.; Schmuki, P., Ta-Doped TiO2 Nanotubes for Enhanced Solar-Light Photoelectrochemical Water Splitting. Chemistry – A European Journal 2013, 19 (19), 5841-5844. 60. Berger, L. I., In CRC Handbook of Chemistry and Physcis,89th ed, Lide, D. R., Ed. CRC Press/Taylor and Francis: Baca Raton, FL, 2008; pp 12(89)-12(77). 61. Turner, J. A., A Realizable Renewable Energy Future. Science 1999, 285, 687-689. 62. Bolton, J. R.; Strickler, S. J.; Connolly, J. S., Limiting and Realizable Efficiencies of Solar Photolysis of Water. Nature 1985, 316, 495-500. 63. Halmann, M., Photoelectrochemical reduction of aqueous carbon dioxide on p-type gallium phosphide in liquid junction solar cells. Nature 1978, 275 (5676), 115-116. 64. McCann, J. F.; Handley, L. J., The photoelectrochemical effect at a p-GaP electrode. Nature 1980, 283 (5750), 843-845. 65. Dare-Edwards, M. P.; Hamnett, A.; Goodenough, J. B., The efficiency of photogeneration of hydrogen at p-type III/V semiconductors. J. Electroanal. Chem. Inter. Electrochem. 1981, 119 (1), 109-123. 66. Erné, B. H.; Vanmaekelbergh, D.; Kelly, J. J., Porous etching: A means to enhance the photoresponse of indirect semiconductors. Adv. Mater. 1995, 7 (8), 739-742. 67. Price, M. J.; Maldonado, S., Macroporous n-GaP in Nonaqueous Regenerative Photoelectrochemical Cells. J. Phys. Chem. C 2009, 113 (28), 11988-11994. 68. Wagner, R. S.; Ellis, W. C., Vapor‐Liquid‐Solid Mechanism of Single Crystal Growth. Applied Physics Letters 1964, 4 (5), 89-90. 69. Lieber, C. M., One-dimensional nanostructures: Chemistry, physics &applications. Solid State Commun. 1998, 107 (11), 607-616. 70. Kayes, B. M.; Filler, M. A.; Putnam, M. C.; Kelzenberg, M. D.; Lewis, N. S.; Atwater, H. A., Growth of vertically aligned Si wire arrays over large areas (>1cm2) with Au and Cu catalysts. Applied Physics Letters 2007, 91 (10), 103110-1-3. 71. Vayssieres, L.; Keis, K.; Lindquist, S. E.; Hagfeldt, A., Purpose-built anisotropic metal oxide material: 3D highly oriented microrod array of ZnO. J. Phys. Chem. B 2001, 105 (17), 3350-3352. 72. Greene, L. E.; Law, M.; Goldberger, J.; Kim, F.; Johnson, J. C.; Zhang, Y. F.; Saykally, R. J.; Yang, P. D., Low-temperature wafer-scale production of ZnO nanowire arrays. Angew. Chem., Int. Ed. 2003, 42 (26), 3031-3034. 73. Fahrenkrug, E.; Gu, J.; Jeon, S.; Veneman, P. A.; Goldman, R. S.; Maldonado, S., Room-Temperature Epitaxial Electrodeposition of Single-Crystalline Germanium Nanowires at the Wafer Scale from an Aqueous Solution. Nano Lett. 2014, 14 (2), 847- 852. 74. Gupta, R.; Xiong, Q.; Mahan, G. D.; Eklund, P. C., Surface Optical Phonons in Gallium Phosphide Nanowires. Nano Lett. 2003, 3, 1745-1750.

75. Xiong, Q.; R.Gupta; Adu, K. W.; Dickey, E. C.; Lian, G. D.; Tham, D.; Fischer, J. E.; Eklund, P. C., Raman Spectroscopy and Structure of Crystalline Gallium Phosphide Nanowires. Journal of Nanoscience and Nanotechnology 2003, 3, 335-339. 76. Seo, H. W.; Bae, S. Y.; Park, J.; Yang, H.; Kang, M.; Kim, S.; Park, J. C.; Lee, S. Y., Nitrogen-Doped Gallium Phosphide Nanobelts. Applied Physics Letters 2003, 82, 3752- 3754. 77. Johansson, J.; Karlsson, L. S.; Svensson, C. P. T.; Mårtensson, T.; Wacaser, B. A.; Deppert, K.; Samuelson, L.; Seifert, W., Structural Properties of <111> B-Oriented III-V Nanowires. Nat. Mater. 2006, 5, 574-580. 78. Dick, K. A.; Caroff, P.; Bolinsson, J.; Messing, M. E.; Johansson, J.; Deppert, K.; Wallenberg, L. R.; Samuelson, L., Control of III-V nanowire crystal structure by growth parameter tuning. Semicond. Sci. Technol. 2010, 25 (2), 024009-1-11. 79. Assali, S.; Zardo, I.; Plissard, S.; Kriegner, D.; Verheijen, M. A.; Bauer, G.; Meijerink, A.; Belabbes, A.; Bechstedt, F.; Haverkort, J. E. M.; Bakkers, E., Direct Band Gap Wurtzite Gallium Phosphide Nanowires. Nano Lett. 2013, 13 (4), 1559-1563. 80. Gu, Z.; Paranthaman, M. P.; Pan, Z., Vapor-Phase Synthesis of Gallium Phosphide Nanowires. Cryst. Growth Des. 2008, 9 (1), 525-527.

CHAPTER II

Photoelectrochemical Behavior of n-type Gallium Phosphide Nanowires as Photoanodes

II.1 Introduction The major challenge in practical PEC is to maximize the photovoltage and the photocurrent. Semiconductors with wide band gaps can produce large photovoltage, but their solar adsorption efficiency is low. Similarly, semiconductors with narrow band gaps can effectively adsorb the solar irradiation, but their photovoltages are limited by the band gap. It is therefore important to choose appropriate materials in order to maximum the efficiency that utilize the solar energy in PEC applications. Gallium phosphide (GaP) has an intrinsic bandgap of 2.26 eV.1 Because of its large photovoltage and relatively high carrier mobility, GaP is a promising material in PEC.2,3,4,5 The bandedges of GaP also align well with several key redox reactions in 3,4,5 6,7,8 energy applications, such as water splitting and CO2 reduction. However, the efficiency of convert solar energy with GaP is limited by its poor light absorption properties. Recent research has shown that this problem can be tackled by using high aspect ratio structures such as macroporous structures9,10,11,12,13 nanowires,14,15,16,17 and microwires.18,19,20,21 The efficiency of high aspect ratio GaP material can reach almost unity from 350 nm up to its bandgap (540 nm).11 The highest performance high aspect ratio GaP photoelectrodes have been prepared through top down (e.g. wafer etching) methods. In order to provide a cost-effective method to prepare high aspect ratio GaP as well as to perform a deep analysis of the photoelectrochemical properties of GaP, in this Chapter we report we report the bottom up synthesis of GaP nanowires film photoelectrode through vapor-liquid-solid (VLS) mechanism (shown in Figure II.1b),22,23 Features such as morphology, crystal defects and catalyst on the nanowire performance

28 were investigated. The photoelectrochemical properties of as-prepared GaP nanowire films were studied in dry acetonitrile solution containing ferrocene/ferrocenium. The critical material properties are identified.

II.2 Methods

II.2.1 Preparation of GaP Nanowires Roughly 15 mgs ground GaP power (99.99% metal basis, Sigma-Aldrich) was placed on a quartz platform. It was then inserted into a quartz tube (diameter 2.2 cm, length 66cm) placed inside a split, single zone tube furnace (Thermolyne F79345). In the meananowirehile, a growth substrate was placed in the same quartz tube about 15 cm away. Single crystalline GaP (100) (miscut ≤ 1o, thickness = 0.5 mm, MTI Corporation) was employed as the growth substrate. The set up is shown in Figure II.1. Before use, each growth substrate was etched in concentrated H2SO4 (doubly distilled, Sigma-Aldrich) for 30 s and rinsed with distilled water. Before growth, the growth substrates were evaporated with a thin (< 4 nm) film of Au prepared by metal evaporation at 4 x 10-6 torr. After the source material and growth substrate were loaded in the tube, the quartz tube was then evacuated to a pressure of 5 x 10-2 torr and left under static vacuum. The quartz tube was heated to 800 °C for 60 min to sublime GaP source powder. After nanowire film deposition, the system was slowly cooled to room temperature over the course of 4 h.

II.2.2 SEM/TEM Characterization of GaP Nanowires GaP nanowire films were sonicated in ethanol for 15s to isolate individual nanowires. These suspensions were then cast onto holey carbon coated TEM copper grids and characterized with a JEOL 3011 transmission electron microscope operated at 300 kV. For collection of electron diffraction patterns, the specimens were tilted and aligned to low index zone axis as determined by the observation of Kikuchi patterns. Scanning electron micrographs of nanowire films were taken separately with a FEI Nova Nanolab SEM/FIB.

II.2.3 Diffuse Reflectance Measurement A 4 inch, four-port Newport integrating sphere coated with Spectraflect® was

Figure II.1 a) A schematic graph of the chemical vapor deposition system used for direct sublimation of source material; b) A schematic graph of Vapor-Liquid-Solid (VLS) nanowire growth, including three steps: I) Alloying; II) Nucleation; III) Axial growth.

Figure II.2 Photoelectrochemical cell (PEC) set up with a flat bottom quartz cell and three electrode system. The tested sample is made as a working electrode.A Platinum coil is used as a counter electrode. A Luggin capillary with platinum inside is the reference electrode.

31 employed for diffuse reflectance measurements of nanowire samples. A 150 W Xeon arc lamp (Newport) coupled to a quarter-turn, single-grating monochrometer (Newport) equipped with 1.24 mm wide rectangular slits was used to produce monochromatic light (± 4 nm tolerance). A Thorlabs S120 Si photodiode detector was coupled to the integrating sphere for collecting reflected light. The chopped incident light was at frequency of 15 Hz and a Stanford Research Systems SR830 lock-in amplifier was used for measuring the photodiode signal. For correcting any variation in lamp intensity, the incident monochromatic was splitted by a quartz beamsplitter to a second Newport 70316NS photodiode detector monitored by a Newport lock-in amplifier (Merlin).The whole system was manipulated by a custom LabVIEW program.

II.2.4 Photoelectrochemistry A glass cell with an optically flat bottom was used for photoelectrochemical measurements. Dry acetonitrile (Aldrich) was obtained from an MBraun solvent purification system. Lithium perchlorate (99.99%, battery grade, Aldrich) was stored and opened in an inert atmosphere glovebox. Ferrocene (Sigma) was sublimed and dried by using schlenk line. Ferrocenium was generated by electrolysis. The cell was assembled in an inert atmosphere glovebox with a platinum as the counter electrode and a luggin capillary reference electrode. After assembly, the cell was removed from the glovebox, and it was kept under a positive pressure of Ar (g) (Metro Welding). J-V characterization curve was collected by CHI Instrument. The absolute photocurrents were measureed with a PAR 273 potentiostat with the working electrode poised at a defined potential. The set up of the photoelectrochemical cell is shown in figure II.2. The ohmic contact on the back side of gallium phosphide wafer was made by soldering a layer of indium on the back and annealed at 400 oC for 10 min with flowing forming gas (5/95 of H2/ N2). KI/I2 solution was prepared by dissolving 1g KI and 0.25g

I2 in 10 ml water. One droplet of KI/I2 solution was placed on the surface of GaP nanowire electrode by different time duration from 1min to 15 min. Photomeasurement was also done in the electrolyte of 1M KCl (Mallinckrodt, Analytical Reagent). In this aqueous electrolyte, a Ag wire coated with AgCl was used as a reference electrode.

II.3 Results

II.3.1 Morphology of GaP Nanowries Figure II.3a shows a cross section SEM image of GaP Nanowires. The thickness of the film was around 20 μm. The prepared Nanowires did not show any preferential alignment. The bottom parts of Nanowires near the substrate were very kinked and twisted. The diameter of Nanowires had a wide distribution range from 100 nm and 400 nm, as indicated in figure II. 3b. The average diameter is around 200 nm. Figure II.3c shows a TEM micrograph of a single GaP nanowire with a diameter around 180 nm. The contrast band along the axis direction of the NANOWIRE in low magnification TEM image indicates the existence of plane defects. By counting the number of bands, the defect concentration is estimated to be around 2 per 100 nm. High resolution TEM image (Figure II.3d) showed single crystalline zincblende lattice. By indexing the diffraction pattern in Fig. II.3d, the growth direction of the nanowire was identified to be [111]. Note that two sets of zincblende pattern at a slight rotational angle were observed in the diffraction pattern, which can be ascribed to the twining boundaries in the lattice.

II.3.2 Photoelectrochemical Response of as-Prepared Nanowire Films Figure II.4 shows a representative J-V response for a GaP nanowire film immersed in solution of ferrocene/ferrocenium in acetonitrile with and without 100 mW·cm-2 white light illumination from ELH lamp. The dark current potential curve shows typical n- type character (inset of figure II.4). The photoresponse shows the evolution of cathodic current when the potential scans towards more negative range. No dark current is observed if the potential scans to more positive values till -0.05 V. When the film is illuminated with white light, the anodic current density was increased up to 0.2 mA·cm-2 at 0 V. By comparing the open circuit under dark and illuminated conditions, the photovoltage of GaP nanowires is estimated to be around 0.8 V. Figure II.5 shows the quantum yield as a function of wavelength for GaP nanowire photoelectrodes. For comparison, the response for GaP planar electrode is also shown. The photoresponse of planar GaP between 350 nm and 550 nm was poor. In this wavelength region, the light absorption is weak. The photon has to penetrate deep into the material to be absorbed and the generated electron/hole pair could not diffuse to the

18 a) 16 14 b) 12 10 8 6

Frequency / % Frequency 4 2 0 100 150 200 250 300 350 400 Diameter / nm c) c) d)

Figure II.3 a) Cross-section SEM image. scale bar: 30 μ m ; b) Diameter distribution of as-prepared GaP nanowire film; c) Low magnification transmission electron microscope image of GaP nanowires, scale bar: 100 nm; Red arrows indicate the positions of twining defects; d) High resolution TEM image of single GaP nanowire, scale bar: 5 nm inset: selected area electron diffraction pattern with [110] beam direction.

-0.3

-0.002

0.000

0.002 -0.2 0.004 0.006 ) 2 0.008

-0.7 -0.6 -0.5 -0.4 -0.3 -0.2 -0.1 Potential (V) -0.1

0.0 illuminated

Current (mA/cm dark 0.1

-0.8 -0.7 -0.6 -0.5 -0.4 -0.3 -0.2 -0.1 0.0 Potential (V)

Figure II.4 J-E behavior of GaP nanowire film in dry acetonitrile with 0.5 mM ferrocenium, 49.5 mM ferrocene and 1M LiClO4. The black curve represents the J-E response in the dark. The red curve represents the J-E characterization under 100 mW·cm-2 of simulated AM 1.5 white light illumination, inset: zoomed in J-E curve in the dark.

0.4 Planar Nanowire-1 Nanowire-2 Nanowire-3 0.3 Nanowire-4 Nanowire-5

0.2

0.1 External Quantum Yield 0.0 350 400 450 500 550 Wavelength ( nm )

Figure II.5 Quantum yield measurement as a function of wavelength of GaP nanowires in dry acetonitrile with 0.5 mM Ferrocenium, 49.5 mM Ferrocene and 1 M LiClO4. The black line represents the photoresponse of p-type GaP crystalline wafer. The other five curves represent five photoelectrodes of GaP nanowire films prepared under the same condition (T=800 oC, t =30 min) in current CVD setup.

Figure II.6 A plot of diffusive reflectance as a function of wavelength of as-prepared GaP nanowire film grown on p-GaP wafer.

37 interface to do useful work before they recombine. The quantum yield of five as-prepared GaP nanowire photoelectrodes was also presented in figure II.5. The quantum yield of GaP nanowires was remarkably improved in the wavelength region from 450 nm to 550 nm. As discussed above, nanowire structure can decouple light absorption and carrier collection direction. In this case, the produced electron/hole pair only needs to diffuse several hundred nanometer to be collected.

II.3.3 Diffusive Reflectance Measurement of GaP Nanowire Films Figure II.6 shows the wavelength dependence of the diffuse reflectance of visible light from a dry GaP nanowire films in air. GaP nanowire film showed minimal reflectance at short wavelengths, in accord with the premise that light with energy greater than the lowest direct bandgap (2.80 eV) of GaP is strongly absorbed by GaP. The comparatively higher diffuse reflectance at wavelengths near and below the lowest indirect bandgap energy (2.26 eV) was consistent with the known visible light-scattering properties of high-aspect ratio GaP morphologies.

II.3.4 Spectral Response of Nanowires in Different Electrolytes The solution of ferrocene and ferrocenium has light yellow and green color, respectively. It suggests the electrolyte adsorbs light in the wavelength region of 350-570 nm. Note this adsorption range coincides with the wavelength for the solar adsorption of GaP. Thus there is a competition between GaP and the electrolyte that adsorbs the solar source. In our experiments, although the photoelectrode has been placed as close as possible to the bottom of the photoelectrochemical cell, it is not avoidable that a layer of solution fills the space between the electrode and the cell. This suggests an optical loss in the electrolyte. We then investigated the quantum yield of GaP Nanowires in different electrolytes. Figure II.7 compares the quantum yield of nanowire film in transparent KCl aqueous solution and yellow colored Fc+/Fc solution. Apparently, the quantum yield in 1 M KCl solution is higher than that in Fc+/Fc nonaqueous solution throughout the wavelength region from 350 nm to 550 nm. Especially in the wavelength region of 350 nm-450 nm, the quantum yield was greatly increased in KCl solution. At wavelength around 365 nm, the quantum yield is 0.079 in Fc+/Fc solution and 0.28 in KCl solution.

0.4

0.3 5 mM Fc/Fc+ 1M KCl

0.2

0.1 External Quantum Yield 0.0 350 400 450 500 550 Wavelength ( nm )

Figure II.7 Quantum yield measurement of GaP nanowire photoelectrodes in two different electrolytes: Ferrocene (49.5 mM) / Ferrocenium (0.5 mM) with 1M LiClO4 in acetonitrile nonqueous solution and 1 M KCl aqueous solution. The black curve represents the photoresponse of GaP nanowire photoelectrode in light yellow nonaqueous electrolyte, the red curve represents the photoresponse of GaP nanowire photoelectrode in transparent electrolyte.

In the wavelength region of 450-550 nm, the quantum yield increases slightly. At the wavelength of 480 nm, the quantum yield is 0.087 in Fc+/Fc solution and 0.097 in KCl solution. Although KCl solution can greatly improve the quantum yield between 350- 450 nm, KI aqueous solution is too corrosive for GaP for long term use.

II.3.5 Effect of Catalyst Residue in the Overall Quantum Yield Performance

The Au catalyst from VLS was removed with KI/I2 etchant. The relevant etching chemistry is believed to be: 24 - - - 2Au + I3 + I →2AuI2 (II.1) According to reference 24, the highest rate to etch planar gold film can reach 1000 Å/minute. For the etching of three-dimensional Au nanostructures, the highest etching rate is ~370 Å/minute.24 This decrease is mainly due to the morphology change - - that slows the diffusion of the large active anions (I3 and AuI2 ) near the surface. Since our material is consisted of three-dimensional nanowire film, it is anticipated that the highest etching rate is even slower. We performed experiments with different etching time to investigate how the etching affects the properties of GaP nanowires. Figure II.7 shows the quantum yield measurement of GaP nanowires etched with

KI/I2. For the sample etched for 1 min, the efficiency in the wavelength region from 350 nm to 400 nm slightly increases. From 400 nm to 500 nm, the quantum yield increases more appreciably. For instance, at 420 nm, the quantum yield was increased from 0.19 to 0.23 after etching. We conclude that removing Au catalyst by a short etching process is beneficial to improve the overall quantum yield. When the nanowire film is etched for 2 min, the quantum yield in the wavelength region from 350 nm to 400 nm noticeably decreased. Especially at wavelength of 370 nm, the quantum yield was decreased from 0.23 to 0.19. However, in the wavelength region of 400 nm- 550 nm, the quantum yield slightly increased. This suggested that longer etching time improves the quantum yield in the wavelength region of 400 nm-500 nm with the trade-off the decreased performance in the wavelength region of 350 nm- 400 nm. For even longer etching time such as 8 min, the quantum yield in the wavelength region of 350 nm- 410 nm is further decreased while the quantum yield reached similar values as the non-etched nanowire film in the wavelength region of 410 nm- 550 nm.

0.30 0 min 1 min 0.25 2 min 8 min 0.20 15 min

0.15 0.10 0.05

External Quantum Yield 0.00 350 400 450 500 550 Wavelength (nm)

Figure II.8 Quantum yield measurement of GaP nanowire photoelectrodes treated with KI/I2 solution for different time duration. The black dashed line represents the photoresponse of as-prepared GaP nanowire photoelectrode. The red line represents the photoresponse of GaP nanowire photoelectrode treated with KI/I2 for 1 min. The blue line represents the photoresponse of GaP nanowire photoelectrode treated with KI/I2 for 15 min.

When the nanowire film is etched for 15 min, the overall quantum yield was noticeably decreased. This effect is more obviously in the wavelength region from 350 nm to 400 nm. For instance, at the wavelength around 365 nm, the quantum yield decreases by 65% after etching. This result suggests that long etching time is negative to the quantum yield. It is probably due to the destruction of the nanowire structure after long etching process. The best etching time in our study is 1 min.

II.4 Discussion From the J-V curve of GaP Nanowires, the as-prepared GaP nanowires show n- type characteristic. This fact suggests the VLS produces n-type GaP nanowires. Several possible reasons can contribute to this observation. First, although we start with stoichiometric GaP powder, the nanowire may grow at a phosphorous rich environment due to the different evaporation and nanowire rate of Ga and P. The enrichment of electron donor leads to n-type GaP. Second, the residue oxygen in the synthesis tube may also be doped in the nanowires, which also donates electrons and results in n-type characteristic. Third, during the growth, Ga and P ratio on the liquid Au alloy and GaP nanowire interface may deviate from the stoichiometric composition. It could also change the composition of the synthesized GaP nanowires. To date, there is one group reported the photoresponse of GaP nanowires prepared by solution synthesis14. Figure II.9 shows the external quantum yield of our nanowires as compared to the reported nanowires. Our nanowires demonstrated much better photoresponse throughout the whole wavelength region (350 nm-550 nm). In addition, from the data in different electrolytes, the yellow electrolyte we were using accounts for appreciable quantum yield loss between 350 and 450 nm. For fair comparison, our prepared nanowire should be expected to have a higher external quantum between 350- 450 nm than the one shown in Figure II.9. By comparing the morphology of these two types of nanowires, our synthesized nanowires (average diameter: 200 nm) shows more straight and much less twisted than the reported nanowires (average diameter: 100 nm). The kinking and diameter sizes are all very influential for the photoelectrochemical performance, which will be further discussed. For all GaP nanowires prepared in this study, the quantum yield is appreciably higher than that for GaP wafer in the wavelength region from 450 to 540 nm. For GaP

0.30 GaP nanowires prepared here 0.25 GaP nanowires reported 0.20 0.15 0.10 0.05

External Quantum Yield 0.00 350 400 450 500 550 Wavelength ( nm )

43 wafer, the adsorption in this region is challenged by the large absorption depth and short distance required to collect photon-induced carriers. Our data demonstrate that this challenge can be tackled by decoupling the photon-adsorption and carrier-collecting with high aspect ratio nanowires. However, in the wavelength region from 350 nm to 450 nm, GaP planar wafer shows better response than nanowire film. We note that our synthesized sample has not been optimized. Thus it is important for future research to systematically optimize the performance of GaP nanowires to improve its carrier collecting properties. Below we discuss several important factors that we propose for further investigations. First, let us discuss the effect of synthetic conditions. In an ideal VLS growth model, the growth of Nanowires follows epitaxial direction of the crystalline substrate. This has been demonstrated in previous study for the VLS growth of Si,25 InP,26 GaP27,28 and other nanowires.29 However, in our study the cross section SEM image suggests the as-prepared nanowires are randomly distributed rather than well orientated along any specific direction. This observation is due to the dense vapor produced in the cold zone of tube furnace, which disturbs the growth of nanowires and results in random growth directions. The SEM image also show a large portion of twisted nanowires, which are resulted from the temperature ramping hysteresis of the single zone furnace system. Therefore it is necessary to precisely control the synthetic conditions such as temperature and/or pressure in order to achieve samples with better qualities in future research. In TEM images, the contrast throughout the axial direction of nanowire indicates the high density of twining planes perpendicular to the growth direction. Twining defects/stacking fault are commonly observed in III-V nanowires30 due to the mixture of wurtzite (WZ) and zinc-blende (ZB) segments. These defects affect the electronic properties of nanowires by scattering electron at twining planes or stacking faults,31,32 which great affect the lifetime and mobility of carriers. This can greatly affect the photoelectrochemical performance.33 Early work suggested that the WZ/ZB segments in III-V nanowires can be monitored by adjusting the temperature and III/V ratio.30,34 In current preparation method, Ga and P vapor are directly sublimed from GaP powder. It is not possible to adjust Ga:P ratio with this method. However, it is possible to add additional Ga or P source in the starting powder to avoid the formation of twinning

44 planes/stacking faults. It is of interest to see how the properties of the nanowires are affected by this method. In our synthesis method, Au was used as the catalyst for the VLS growth of GaP nanowires. During the growth process, it is possible that certain amount of Au may also incorporate into the nanowires or wet the side wall of the nanowires.35,36 Because gold is a well-known deep-level trap for semiconductors,37 removing gold from the nanowires is essential to improve electrical properties of nanowires. Our KI etching experiments show that by etching one minute in KI solution, the quantum yield can be noticeably improved. However, with longer etching time, the efficiency is greatly decreased. Longer time etching can roughen the surface of nanowire and be detrimental to the nanowire structure. This could increase surface recombination rate and decrease the efficiency. Finally, let us discuss further about the physical origin of the lower quantum yield in the wavelength region from 350 to 450 nm for GaP nanowires. In previous simulation work,38 it was shown that the relative length scale between radius (r) of the nanowire and depletion width (W) of the semiconductor is a key descriptor for high-aspect-ratio heterojunctions. Nanowires with r < W cannot support the full electric field between the surface and bulk regions. In this case the majority carriers cannot be driven away from the interface, resulting in recombination losses at the interface. This decreases the attainable quantum yield of charge carriers. This explains as-prepared nanowires did show better response than the single crystalline wafer, but the 350-450 nm photoresponse was worse than single crystalline wafer. To improve the quantum yield of nanowires, it is necessary to obtain nanowires with r > W. One strategy to achieve this target is to use nanowires with larger diameters. However, it would eventually leads to the usage of bulk GaP, or GaP wafers, which has low efficiency at high wavelength. Another alternative strategy is to decrease W of GaP. As shown in Eq. II.2, the depletion width is determined as

2  W  0 b (II.2) qN D where ε is the effective dielectric constant of the semiconductor, ε0 is the permittivity of free space, Φb is the equilibrium junction barrier height, q is the unsigned charge of an electron, ND is the doping density.

An effective strategy to decrease W is to increase the doping density. This will be discussed further in Chapter IV.

II.5 Conclusions The photoelectrochemical conversion efficiency of GaP in the solution of ferrocene/ferrocenium in acetonitrile was improved in the wavelength region from 450 nm to 550 nm by using nanowires prepared from low-grade materials. The further improvement of the nanowires can be achieved by the brief etching in KI/I2 solution. The observation of relative low efficiency in the wavelength region from 350 nm to 550 nm can be attributed to non-effective doping density. The data suggests the use of nanowires prepared from low-grade material can be one effective strategy to use GaP as useful material for photoelectrochemical application.

II.6 References 1. Strehlow, W. H.; Cook, E. L., Compilation of Energy Band Gaps in Elemental and Binary Compound Semiconductors and Insulators. J. Phys. Chem. Ref. Data 1973, 2, 163-199. 2. Kohl, P. A.; Bard, A. J., Semiconductor electrodes. 13. Characterization and behavior of n-type zinc oxide, cadmium sulfide, and gallium phosphide electrodes in acetonitrile solutions. J. Am. Chem. Soc. 1977, 99 (23), 7531-7539. 3. McCann, J. F.; Handley, L. J., The photoelectrochemical effect at a p-GaP electrode. Nature 1980, 283 (5750), 843-845. 4. Dare-Edwards, M. P.; Hamnett, A.; Goodenough, J. B., The efficiency of photogeneration of hydrogen at p-type III/V semiconductors. J. Electroanal. Chem. Inter. Electrochem. 1981, 119 (1), 109-123. 5. Ohashi, K.; McCann, J.; Bockris, J. O. M., Stable Photoelectrochemical Cells for Splitting of Water. Nature 1977, 266 (5603), 610-611. 6. Halmann, M., Photoelectrochemical reduction of aqueous carbon dioxide on p-type gallium phosphide in liquid junction solar cells. Nature 1978, 275 (5676), 115-116. 7. Petit, J.-P.; Chartier, P.; Beley, M.; Deville, J.-P., Molecular catalysts in photoelectrochemical cells: Study of an efficient system for the selective photoelectroreduction of CO2: p-GaP or p-GaAs/Ni(cyclam)2+, aqueous medium. J. Electroanal. Chem. Inter. Electrochem. 1989, 269 (2), 267-281. 8. Barton, E. E.; Rampulla, D. M.; Bocarsly, A. B., Selective Solar-Driven Reduction of CO2 to Methanol Using a Catalyzed p-GaP Based Photoelectrochemical Cell. J. Am. Chem. Soc. 2008, 130 (20), 6342-6344. 9. Sivula, K.; Zboril, R.; Le Formal, F.; Robert, R.; Weidenkaff, A.; Tucek, J.; Frydrych, J.; Grätzel, M., Photoelectrochemical Water Splitting with Mesoporous Hematite Prepared by a Solution-Based Colloidal Approach. J. Am. Chem. Soc. 2010, 132 (21), 7436-7444.

10. Maiolo, J. R.; Atwater, H. A.; Lewis, N. S., Macroporous Silicon as a Model for Silicon Wire Array Solar Cells. The Journal of Physical Chemistry C 2008, 112 (15), 6194-6201. 11. Price, M. J.; Maldonado, S., Macroporous n-GaP in Nonaqueous Regerative Photoelectrochemical Cells. 2009, 113 (28), 11988-11994. 12. Hagedorn, K.; Collins, S.; Maldonado, S., Preparation and Photoelectrochemical Activity of Macroporous p-GaP (100). J. Electrochem. Soc. 2010, 157 (11), D588-D592. 13. Brillet, J.; Grätzel, M.; Sivula, K., Decoupling Feature Size and Functionality in Solution-Processed, Porous Hematite Electrodes for Solar Water Splitting. Nano Lett. 2010, 10 (10), 4155-4160. 14. Liu, C.; Sun, J. W.; Tang, J. Y.; Yang, P. D., Zn-Doped p-Type Gallium Phosphide Nanowire Photocathodes from a Surfactant-Free Solution Synthesis. Nano Lett. 2012, 12 (10), 5407-5411. 15. Goodey, A. P.; Eichfeld, S. M.; Lew, K.-K.; Redwing, J. M.; Mallouk, T. E., Silicon nanowire array photoelectrochemical cells. J. Am. Chem. Soc. 2007, 129 (41), 12344- 12345. 16. Yuan, G.; Zhao, H.; Liu, X.; Hasanali, Z. S.; Zou, Y.; Levine, A.; Wang, D., Synthesis and Photoelectrochemical Study of Vertically Aligned Silicon Nanowire Arrays. Angew. Chem., Int. Ed. 2009, 48 (51), 9680-9684. 17. Dalchiele, E. A.; Martín, F.; Leinen, D.; Marotti, R. E.; Ramos-Barrado, J. R., Single- Crystalline Silicon Nanowire Array-Based Photoelectrochemical Cells. J. Electrochem. Soc. 2009, 156 (5), K77-K81. 18. Boettcher, S. W.; Warren, E. L.; Putnam, M. C.; Santori, E. A.; Turner-Evans, D.; Kelzenberg, M. D.; Walter, M. G.; McKone, J. R.; Brunschwig, B. S.; Atwater, H. A.; Lewis, N. S., Photoelectrochemical Hydrogen Evolution Using Si Microwire Arrays. J. Am. Chem. Soc. 2011, 133 (5), 1216-1219. 19. Strandwitz, N. C.; Turner-Evans, D. B.; Tamboli, A. C.; Chen, C. T.; Atwater, H. A.; Lewis, N. S., Photoelectrochemical Behavior of Planar and Microwire-Array Si|GaP Electrodes. Advanced Energy Materials 2012, 2 (9), 1109-1116. 20. Santori, E. A.; Maiolo, J. R., III; Bierman, M. J.; Strandwitz, N. C.; Kelzenberg, M. D.; Brunschwig, B. S.; Atwater, H. A.; Lewis, N. S., Photoanodic behavior of vapor- liquid-solid-grown, lightly doped, crystalline Si microwire arrays. Energy Environ. Sci. 2012, 5 (5), 6867-6871. 21. Mallorqui, A. D.; Epple, F. M.; Fan, D.; Demichel, O.; Morral, A. F. I., Effect of the pn junction engineering on Si microwire-array solar cells. Physica Status Solidi a- Applications and Materials Science 2012, 209 (8), 1588-1591. 22. Wagner, R. S.; Ellis, W. C., Vapor‐Liquid‐Solid Mechanism of Single Crystal Growth. Applied Physics Letters 1964, 4 (5), 89-90. 23. Wu, Y.; Yang, P., Direct Observation of Vapor−Liquid−Solid Nanowire Growth. J. Am. Chem. Soc. 2001, 123 (13), 3165-3166. 24. T. Hattori, J. R., R. Novak, P. Merten, P.Besson, Cleaning and Surface Conditioning Technology in Semiconductor Device Manufacturing 11. The Electrochemical Society: 2009; p 393. 25. Kayes, B. M.; Filler, M. A.; Putnam, M. C.; Kelzenberg, M. D.; Lewis, N. S.; Atwater, H. A., Growth of vertically aligned Si wire arrays over large areas (>1cm2) with Au and Cu catalysts. Applied Physics Letters 2007, 91 (10), 103110-1-3.

26. Mårtensson, T.; Carlberg, P.; Borgström, M.; Montelius, L.; Seifert, W.; Samuelson, L., Nanowire Arrays Defined by Nanoimprint Lithography. Nano Lett. 2004, 4 (4), 699- 702. 27. Borgstrom, M. T.; Immink, G.; Ketelaars, B.; Algra, R.; Bakkers, E., Synergetic nanowire growth. Nature Nanotechnology 2007, 2 (9), 541-544. 28. Assali, S.; Zardo, I.; Plissard, S.; Kriegner, D.; Verheijen, M. A.; Bauer, G.; Meijerink, A.; Belabbes, A.; Bechstedt, F.; Haverkort, J. E. M.; Bakkers, E., Direct Band Gap Wurtzite Gallium Phosphide Nanowires. Nano Lett. 2013, 13 (4), 1559-1563. 29. Tomioka, K.; Motohisa, J.; Hara, S.; Fukui, T., Control of InAs Nanowire Growth Directions on Si. Nano Lett. 2008, 8 (10), 3475-3480. 30. Caroff, P.; Dick, K. A.; Johansson, J.; Messing, M. E.; Deppert, K.; Samuelson, L., Controlled polytypic and twin-plane superlattices in III-V nanowires. Nature Nanotechnology 2009, 4 (1), 50-55. 31. Stiles, M. D.; Hamann, D. R., Ballistic Electron Transmission Through Interfaces. Phys. Rev. B 1988, 38 (3), 2021-2037. 32. Stiles, M. D.; Hamann, D. R., Electron Transmission Through Silicon Stacking-Faults. Phys. Rev. B 1990, 41 (8), 5280-5282. 33. Foley, J. M.; Price, M. J.; Feldblyum, J. I.; Maldonado, S., Analysis of the operation of thin nanowire photoelectrodes for solar energy conversion. Energy Environ. Sci. 2012, 5 (1), 5203-5220. 34. Dick, K. A.; Caroff, P.; Bolinsson, J.; Messing, M. E.; Johansson, J.; Deppert, K.; Wallenberg, L. R.; Samuelson, L., Control of III-V nanowire crystal structure by growth parameter tuning. Semicond. Sci. Technol. 2010, 25 (2), 024009-1-11. 35. Allen, J. E.; Hemesath, E. R.; Perea, D. E.; Lensch-Falk, J. L.; Li, Z. Y.; Yin, F.; Gass, M. H.; Wang, P.; Bleloch, A. L.; Palmer, R. E.; Lauhon, L. J., High-resolution detection of Au catalyst atoms in Si nanowires. Nature Nanotechnology 2008, 3 (3), 168-173. 36. (a) Hannon, J. B.; Kodambaka, S.; Ross, F. M.; Tromp, R. M., The influence of the surface migration of gold on the growth of silicon nanowires. Nature 2006, 440 (7080), 69-71; (b) Koren, E.; Elias, G.; Boag, A.; Hemesath, E. R.; Lauhon, L. J.; Rosenanowireaks, Y., Direct Measurement of Individual Deep Traps in Single Silicon Nanowires. Nano Lett. 2011, 11 (6), 2499-2502. 37. Collins, C. B.; Carlson, R. O.; Gallagher, C. J., Properties of Gold-Doped Silicon. Physical Review 1957, 105 (4), 1168-1173. 38. Hagedorn, K.; Forgacs, C.; Collins, S.; Maldonado, S., Design Considerations for Nanowire Heterojunctions in Solar Energy Conversion/Storage Applications. J. Phys. Chem. C 2010, 114 (27), 12010-12017.

CHAPTER III

Structural and Photoelectrochemical Properties of Gallium Phosphide Nanowires

Annealed in NH3

*This work is adapted from a published work Wen, W.; Carim, A.; Collins, S.; Price, M.; Peczonczyk, S.; Maldonado, S.; J.Phys.Chem.C ,2011, 115,22652

III.1 Introduction New semiconductor photoelectrode materials are needed to advance artificial photosynthetic systems1 for solar energy conversion/storage. For photoelectrochemical cells that store incident sunlight in energy-rich chemical fuels, candidate materials should demonstrate two important properties. First, the photoelectrode material should absorb and convert an appreciable portion of the solar spectrum. This constraint implies that the semiconductor photoelectrode passes a relatively high photocurrent density under illumination at the level of 1.5 AMU.2 Second, the thermodynamics of fuel-forming reactions require that the semiconductor photoelectrode generates a cell electromotive force greater than the standard potential of the reaction. For solar water splitting, the cell voltage requirement is 1.23 V at standard conditions.3 The mechanistically complex processes for bond formation/scission in water splitting demand an additional overpotential that increases the total needed photo-induced voltage to ≥ 1.7 V.4,5 Mid- sized bandgap semiconductors are naturally suited to simultaneously satisfy both requirements and are thus appropriate for solar-powered fuel generation systems.6 Gallium phosphide (GaP) has long been considered as a candidate semiconductor for photosynthetic photoelectrochemical applications since its mid-sized bandgap intrinsically allows for large photovoltages under illumination.7 Unfortunately, short minority carrier diffusion lengths typical in GaP artificially lower the capacity for

49 collection of long wavelength light by planar GaP photoelectrodes. However, two separate strategies have been explored to increase the attainable photocurrent density with GaP photoelectrodes. Our recent efforts have demonstrated that GaP photoelectrodes with short minority carrier diffusion lengths can exhibit excellent photoresponse characteristics for photon energies as low as the bandgap energy if a high aspect ratio form factor is used.8,9 Separate work by Turner et. al. has shown that the bandgap energy of GaP can be lowered slightly through alloying GaP with N to form 10 GaP1-xNx (0 ≤ x ≤ 1). This work also demonstrated the excellent corrosion resistance of 10 GaP1-xNx in aqueous electrolytes, an additional feature attractive for practical water electrolysis systems. To date, GaP photoelectrodes that have both a high aspect ratio and an appreciable level of N have not been examined as possible solar energy conversion/storage materials. The advantage of combining these two strategies is that the high-aspect ratio form factor increases the capacity for efficient light energy capture/conversion out to the bandgap wavelength and that N alloying lowers the energy threshold for optical energy capture and conversion. Long nanowire films naturally embody one type of high-aspect ratio form factor and have been accordingly studied as a viable photoelectrode design for widescale use.11,12 Herein, we show data illustrating that

GaP1-xNx nanowire films can be prepared by annealing GaP nanowire films in NH3(g) under conditions that do not result in a total conversion to gallium nitride (GaN).13 In contrast to a previous report on GaP nanowires with only an outer shell enriched with N,14 we present data illustrating N incorporation throughout the nanowire volume. In addition, we report data detailing the physicochemical and photoelectrochemical properties of GaP1-xNx nanowire films. In the context of Raman scattering as a probe for the local structural order in III-V semiconductors following ion irradiation,15 doping,16,17 and alloying with either isoelectronic cationic or anionic substitutional impurities,18 we report and interpret the rich and unusual features of the Raman spectra of GaP nanowires following treatment with NH3(g). The sum materials analyses are used to assess this methodology for producing alloyed GaP nanowire films and the prospects for preparing high efficiency, GaP-based photoelectrodes capable of sustaining large photovoltages and high photocurrents are discussed.

III.2 Methods

III.2.1 GaP Nanowire Film Preparation A small amount (ca. 15 mg) of ground GaP power (99.99% metal basis, Sigma- Aldrich) was placed on a quartz platform which was then inserted into a quartz tube (diameter 2.2 cm, length 66cm) placed inside a split, single zone tube furnace (Thermolyne F79345). Concomitantly, a growth substrate was loaded into the same quartz tube. Two types of growth substrates were employed, single crystalline Si and single crystalline GaP (111A) (miscut ≤ 0.5º, thickness = 0.5 mm, MTI Corporation). Before use, each growth substrate was etched, with the Si substrates etched in hydrofluoric acid (49% (aq) by weight) for 30 s and rinsed with distilled water (>18 MΩ cm, Barnstead Nanopure III purifier) and with GaP sections etched in concentrated

H2SO4 (doubly distilled, Sigma-Aldrich) for 30 s and rinsed with distilled water. Prior to loading, the growth substrates were coated with a thin (< 4 nm) film of Au prepared by metal evaporation at 4 x 10-6 torr. Upon loading the source material and growth substrate, the quartz tube was evacuated to a pressure of 5 x 10-2 torr and left under static vacuum. The quartz tube was then heated to 800 °C for 1 h to sublime the solid GaP source material. After nanowire film deposition, the system was slowly cooled to room temperature over the course of 4 h. For N alloying, the as-prepared GaP nanowire films were subsequently annealed in a quartz flow tube system (diameter 4.4 cm, length 90 cm) at temperatures ranging between 600 and 750 °C in flowing NH3(g). The system was initially purged of O2(g) by flowing Ar(g) (Cryogenic Gases) at 450 sccm for 30 min at room temperature. The temperature was then raised to the desired setpoint temperature and NH3(g) (Cryogenic

Gases) was introduced into the gas stream. The total flow rate of Ar(g) and NH3(g) was kept at 150 sccm for 4 h. The ratio of NH3(g):Ar(g) was used to control the level of nitrogen incorporation. Upon completion, the NH3(g) flow was stopped and the system was allowed to cool radiatively to room temperature over a time period of 4 h.

III.2.2 Materials Characterization Samples for transmission electron microscopy analysis were prepared by first sonicating GaP nanowire films in ethanol for 15 s to dislodge individual nanowires from

51 the growth substrate. These suspensions were then cast onto holey carbon coated TEM copper grids and characterized with a JEOL 3011 transmission electron microscope operated at 300 kV. For collection of electron diffraction patterns, the specimens were tilted and aligned to low index zone axis as determined by the observation of Kikuchi patterns. Scanning electron micrographs of nanowire films were taken separately with a FEI/Philips XL30 FEG. X-ray diffraction data were acquired with a Rigaku Ultima IV diffractometer equipped with a Cu Kα source, a steel sample support plate, and a scintillation counter detector. A grazing angle geometry was employed to eliminate the contribution of reflections from the underlying substrate. The sample position and incident angle were first aligned by monitoring the transmission of the incident X-rays with respect to the stage height and incident angle and then precisely adjusted using the reflected X-rays so that the beam exactly entered the receiving slit. During pattern collection, the incoming X-ray beam was fixed at an incident angle of 0.3° with respect to the sample plane while the detector was swept through the full 2θ range. Diffuse reflectance measurements of representative nanowire samples annealed at

750°C under various NH3-containing ambients were taken using a 4 inch, four-port Newport integrating sphere coated with Spectraflect®. Monochromatic light (± 4 nm tolerance) was generated from the output of a 150 W Xe arc lamp (Newport) coupled to a quarter-turn, single-grating monochrometer (Newport) equipped with 1.24 mm wide rectangular slits. Reflected light was collected with a a Thorlabs S120 Si photodiode detector coupled to the integrating sphere. The incident light was chopped at 15 Hz and the photodiode signal was measured with a Stanford Research Systems SR830 lock-in amplifier. A quartz beamsplitter was used to direct a portion of the incident monochromatic beam to a second Newport 70316NS photodiode detector monitored by a Newport lock-in amplifier (Merlin) to correct for variation in lamp intensity. The system was controlled via a custom LabVIEW program. Raman spectra of as-prepared and nitrogen-alloyed GaP nanowire films on Si substrates were obtained with a Renishaw inVia Raman spectrometer equipped with a Leica microscope, an Olympus SLMPlan 20x objective (numerical aperture = 0.35), holographic notch filter, a 1800 lines/mm grating and a RenCam CCD detector in a 180°

52 backscatter geometry. A 514.5 nm Ar+ laser (Laser Physics 25s) and a 632.8 nm HeNe laser (Renishaw RL633) were used as excitation sources with radiant fluxes of 500 μW and 390 μW, respectively, at the sample. The incident excitation was linearly polarized and no polarizing collection optic was used for spectral acquisition. For single nanowires, polarization optics was used to collect Raman spectra at 514.5 nm excitation. Dispersions of single nanowires were prepared by sonicating nanowire films in ethanol and drop- casting the resultant suspensions onto glass microscope slide sections. A Renishaw inVia Raman spectrometer equipped with a Leica microscope, a Leica N Plan 100x objective (numerical aperture = 0.90), holographic notch filter, a 1800 lines/mm grating and a RenCam CCD detector was employed. A 514.5 nm Ar+ laser (Laser Physics 25s) was used as the excitation source with a radiant flux of 500 μW at the sample. The system was configured in 180˚ backscatter geometry in which the incident excitation and collected Raman scatter propagated along the Y axis (surface normal). The wires, depicted in Figure III.1, were oriented with their long axes along the Z axis, perpendicular to the intrinsic polarization direction of the laser (X axis). A half-wave plate was placed in the incident beam path to rotate the polarization direction parallel to the long axes (Z axis) when necessary. Meanwhile, a linear polarizer was placed between the notch filter and the grating and rotated to isolate Raman scatter with polarization along either the Z or X axes for detection. The convention used to define the polarization condition for each spectra acquisition is diagrammed in Figure III.1. X-ray photoelectron (XP) spectra were acquired with a PHI 5400 analyzer using a Mg Kα (1253.6 eV) source without a monochromator at a takeoff angle of 54.6°. Spectra were recorded without charge neutralization at a base pressure of < 2.0×10-9 Torr. A 6 mA emission current and a 10 kV anode high tension (HT) were used. The binding energies (BEs) of all spectra were corrected using the expected BE for adventitious carbon (284.6 eV). Typical acquisition times for high-resolution spectra ranged from 30 to 45 min. The obtained XP spectra were analyzed using CASAXPS software Version 2.313. A commercially available crystalline GaN film (undoped GaN template on c-plane sapphire, Kyma Technologies) was used to collect a reference N1s spectrum.

III.2.3 Photoelectrochemical Measurements Photoelectrode Preparation For photoelectrochemical measurements, ohmic contacts to the nanowire film electrodes were prepared by etching the back side of each substrate in H2SO4 for 30 s, rinsing with distilled water, soldering a thin film of In (99.9+% metal basis, Aldrich) on the back, and then annealing in forming gas (95:5

N2:H2) at 400°C for 10 min. Electrodes were then prepared by attaching the back side of each nanowire film to a copper wire coil threaded through a glass tube with silver epoxy (GC Electronics) and sealing with inert epoxy (Hysol C). A glass cell with an optically flat bottom serving as the window was used for photoelectrochemical measurements. Dry acetonitrile (Aldrich) was prepared with an MBraun solvent purification system. Lithium perchlorate (99.99%, battery grade, Aldrich) was opened in an inert atmosphere glovebox, and used as received. Ferrocene (Sigma) was sublimed and dried before use, and ferrocenium was generated electrolytically with a second compartment separated by a Vycor frit. The cell was assembled in an inert atmosphere glovebox and included a platinum gauze counter electrode and a luggin capillary reference electrode containing a platinum wire poised at the solution potential. After assembly, the cell was removed from the glovebox, connected to an Ar (g) line, and kept under a positive pressure of Ar (g) (Metro Welding) during use. Quantum efficiency measurements were obtained using the spectral response system described above. The absolute photocurrents were measured with a PAR 273 potentiostat with the working electrode poised at short-circuit.

III.3 Results

III.3.1 X-ray Diffraction Figure III.2a shows a representative electron micrograph from a cross-sectional viewpoint of a GaP nanowire film deposited on a planar substrate. Each investigated film in this report was composed of crystalline GaP nanowires with nominal diameters of 80 nm and lengths in excess of 20 μm. Annealing GaP nanowires at 800 °C for 4 hrs in pure flowing NH3(g) did not elicit any perceptible change in the nanowire film morphology (i.e. height, density), although roughening of the nanowire surfaces was evident in some scanning electron micrographs. X-ray diffractograms collected for GaP nanowire films

GaP GaN a) b) (200) (002) + GaN GaP GaP GaN (101) (111) (400) (100) + GaP GaN GaN Fe (220) (112) (201) (110) GaP (311) GaP GaN (331)

Intensity (a. u.) (110)

20 40 60 80 2020μmm 2 (o)

GaP GaP (200) c) GaN GaP d) (111) + GaN (002) (101) (400) GaN + Fe GaP (100) GaN (110) (111) (112) GaP GaP GaP (220) GaP GaP (220) Fe (311) (331) (200) GaP GaP (110) GaN (311) GaP (331) (110) (400) Intensity (a. u.) Intensity (a. u.)

20 40 60 80 20 40 60 80 o o 2( ) 2( ) Figure III.2 (a) Cross sectional view of a scanning electron micrograph of a GaP nanowire film grown on Si (100) substrate. (b-d) X-ray diffractograms of GaP nanowire films annealed in (b) pure NH3(g) at T = 800 °C, (c) pure NH3(g) at T = 750 °C, and (d) NH3(g):Ar(g) (1:5) at T = 800 °C.

56 annealed under these conditions showed crystalline reflections for both zincblende GaP and wurtzite GaN (Figure III.2b), indicating enough N incorporation to effect a complete conversion from GaP to GaN in at least some portions of the films. Diluting the NH3(g) in the effluent gas at this temperature did not reproducibly prevent the undesired formation of GaN. To minimize the propensity for conversion of GaP to GaN under annealing in NH3(g), lower annealing temperatures were explored. At temperatures below 700 °C, the alloying reaction was slow and difficult to reproduce. At 750 °C, the conversion of GaP to GaN was still observed when pure NH3(g) was used (Figure III.2c).

However, diluting the NH3(g) content in the gas stream at this temperature showed different results. Figure III.2d shows a representative X-ray diffractogram with no detectable GaN-based reflections from the annealed nanowire film after heating at 750 °C in a dilute NH3(g) (1:5 in Ar(g)) flowing gas stream for 4 hrs. Repeated measurements of nanowire films annealed in even more dilute NH3(g) also showed no diffraction-based evidence of GaN.

III.3.2 Diffuse Reflectance Although no perceptible changes in the X-ray diffractograms were noted for GaP nanowire films after annealing at 750 °C in flowing pure and dilute NH3(g), the tint of these films changed noticeably. Pristine GaP nanowire films with features as shown in Figure III.3a had a different hue than polished GaP wafers, consistent with previous observations by Muskens et. al.20 These nanowire films natively possessed a dull, pale- yellow tint (Figure III.3a). Annealing in the most dilute NH3(g):Ar(g) ambient (1:20 v/v) produced a change in the color of GaP nanowire films, shifting to a deeper yellow.

Analogous annealing steps in more concentrated NH3(g) gas mixtures caused a steady change from yellow to a deep orange-red color (Figure III.3a). Figure III.3b shows the wavelength dependence of the diffuse reflectance of visible light from a dry GaP nanowire films in air after annealing in several different NH3(g) atmospheres at 750 °C. All GaP nanowire films showed minimal reflectance at short wavelengths, in accord with the premise that light with energy greater than the lowest direct bandgap (2.76 eV) of GaP is strongly absorbed by GaP.7 For the pristine GaP nanowire film, the comparatively higher diffuse reflectance at wavelengths near and below the lowest indirect bandgap energy was consistent with the known visible light-scattering properties of high-aspect

1.0 NH :Ar (1:9) a) 3 b) No Anneal NH :Ar (1:20) 0.8 3 NH3:Ar (1:14) NH :Ar (1:5) NH :Ar (1:9) No Anneal 3 3 NH3:Ar (1:5)

0.6 NH3:Ar (1:1)

NH3:Ar (1:0)

NH3:Ar (1:20) NH3:Ar (1:1) 0.4

NH :Ar (1:14) NH :Ar (1:0) 3 3 0.2 Diffuse Reflectance

0.0 300 400 500 600 700 Wavelength (nm) Figure III.3 (a) Plan view optical images of GaP nanowire films as shown in Figure III.2a after annealing treatments at T = 750 °C with various concentrations of NH3(g). (b) Spectral profiles of the diffuse reflectance of light collected from the top of dry GaP nanowire films in air. The designations denote the annealing conditions for each sample, as in Figure III.3a.

21 ratio GaP morphologies. For GaP nanowire films annealed in NH3(g), the reflectance measured at short wavelengths (< 450 nm) remained nominally unchanged. Lower light reflectance was evident at wavelengths between 450 and 650 nm for all GaP nanowire films annealed in NH3(g), with lower reflectance values noted for GaP nanowire films annealed with higher concentrations of NH3(g).

III.3.3 Raman Spectroscopy Raman spectra were collected to better ascertain the nature of structural changes induced by N incorporation. Figure III.4a summarizes the Raman spectra for GaP nanowire films collected with a backscattering orientation, excitation at λ = 514.5 nm, and no polarization optics for either excitation or collection. Between 100-600 cm-1, the Raman spectra for pristine GaP nanowires showed two prominent signatures at 365 and 403 cm-1. The peak positions of these two features were consistent with GaP lattice phonons arising from the Brillouin zone center (Γ point) and thus were assigned as the transverse optical (TO(Γ)) and longitudinal optical (LO(Γ)) phonon modes.22 A low frequency shoulder on the LO(Γ) mode was also observed, consistent with a surface optical phonon mode.22c One additional weak signature was also noted at 210 cm-1, the expected frequency for the first overtone of the transverse acoustic mode at the X point in the Brillouin zone (TA(X)).22d GaP nanowire films annealed at T = 750 °C in Ar(g) without NH3(g) showed no significant change in relative intensity, occurrence, or width in any of these spectral features (Figure III.4), indicating annealing at this temperature did not significantly distort the zincblende GaP lattice. In contrast, GaP nanowire films annealed at T = 750 °C and in the presence of NH3(g) did exhibit pronounced changes in their Raman spectra. Figure III.4a summarizes the Raman spectra for GaP nanowire films treated as described in Figure III.3. These same data are reproduced as Figures III.5 (a)-(d) which separately focus on four different spectral ranges for further comparison. Additionally, the assignments for each observed signature for all of the spectra in Figure

III.4 are summarized in Table III.1. The Raman spectra collected for the NH3(g)-treated GaP nanowire films still showed the same principal TA(X), TO(Γ), and LO(Γ) modes with only minor differences. The relative intensity of the TA(X) overtone noticeably increased and the frequency of the LO(Γ) mode red-shifted slightly after annealing in the presence of NH3(g) (Table 1). These spectra also contained several additional signatures.

a)1:1 20 x b) c) Y(ZX)Y 15 x 5x 1:5 20 x Y(ZX)Y 5x

5x 1:9 20 x Y(ZZ)Y Y(ZZ)Y 20 x 1:14 20 x Y(XZ)Y 1:20 2x Y(XZ)Y

Intensity (a.u.) 4x 20 x 0:1 Y(XX)Y 20 x 10 x No annealing Y(XX)Y 4x

200 400 600 100 200 300 400 500 600 100 200 300 400 500 600 -1 -1 -1 Raman Shift (cm ) Raman Shift (cm ) Raman Shift (cm ) Figure III.4 (a) Representative Raman spectra for GaP nanowire films collected with excitation at 514.5 nm. The spectra are offset for clarity and normalized to the intensity of the phonon mode at 365 cm-1. The insets correspond to the same spectral data magnified by a factor of 20. Spectra were collected without any excitation or collection polarizing optics. (b) Polarized Raman spectra of an individual GaP nanowire cast on a glass slide. Light was incident and collected at the surface normal (Y vector). The terms in the parentheses indicate the polarization direction of the incident and collected light, respectively. For these measurements, the long axis of the nanowire was parallel to the Z- axis. The spectra are offset vertically for clarity and have been rescaled by the designated factor for ease of comparison. (c) Same as in (b) except for a GaP nanowire after annealing in a gas stream of NH3(g) and Ar(g) (1:5, v/v).

The Raman spectra showed 11 new prominent spectral features (Table III.1) with intensities that directly tracked the relative proportion of NH3(g) used during the annealing step. Seven of the new signatures occurred at frequencies commensurate with linear combinations of GaP phonon modes arising from the X and L points of the Brillouin zone.23 The increased relative intensity of the TA(X) signature also tracked the

NH3(g) content used during annealing. Three additional Raman modes appeared as low frequency shoulders on the TO(Γ) and LO(Γ) modes at frequencies consistent with assignments of TO(X)/LO(X), LO(L), and LO(X) phonon modes, respectively. The last new feature common to all of the NH3(g)-annealed GaP nanowire samples was observed at 431 cm-1 but did not directly correspond to any first or second-order zincblende lattice phonons at the X or L points of the Brillouin zone of GaP.23 A possible assignment could be as the overtone of the AIII zone edge phonon at the critical point W in the Brillouin zone23. A separate mode at 550 cm-1, consistent with the TO(Γ) phonon of wurtzite 24 GaN, was only observed for GaP nanowire samples treated with 1:1 v/v NH3(g):Ar(g) or in pure NH3(g). This same Raman signature was even more pronounced for spectra collected for GaP nanowire films annealed at T > 750 ºC which also showed X-ray crystallographic evidence of GaN (Figure III.2). To identify resonant modes from modes induced by nitrogen incorporation, Raman spectra with different excitation wavelengths (514.5 nm Ar+ laser and 632.8 nm HeNe laser) were collected. Figure III.6 illustrates resonant effects on the GaP nanowires annealed in a gas stream of NH3 (g) and Ar (g) (1:5, v/v) at 750 °C. Only two of the new features (modes at 279 and 302 cm-1) in the

Raman spectra for the GaP nanowire films annealed in NH3(g) showed any evidence of resonant enhancements between excitation at 633 nm (1.96 eV) and 514 nm (2.41 eV). Additional annealing in Ar(g) for 2 h at T = 750 ºC did not cause any perceptible change in the relative intensities or line widths observed in the Raman spectra. Polarized Raman spectra were also collected for individual GaP nanowires. Figure III.4b summarizes the observed Raman spectra collected for a pristine GaP nanowire. The intensity of the LO(Γ) mode relative to the intensity of the TO(Γ) mode was strongly sensitive to the polarization conditions, in accord with separate reports from Chapelle et. al.22d and Wu et. al.25 Figure III.4c shows polarized Raman spectra for a single GaP nanowire that had been annealed in 1:5 v/v NH3(g):Ar(g) at T = 750 ºC. In

Figure III.5 (a)-(d) show Raman spectra of representative gallium phosphide nanowire films highlighted in different wavenumber ranges.

Table III.1 Raman spectral features for GaP nanowires annealed at T = 750°C.

Peak Position /cm-1 Assignments Relative Peak Intensitya

NH3(g):Ar(g) (v/v) (0:1) (1:20) (1:14) (1:9) (1:5) (1:1)

178 2TA(L) - 0.00043 0.011 0.050 0.070 0.20

209 2TA(X) 0.022 0.023 0.032 0.050 0.050 0.15

250 TO(X)-TA(X) - 0.011 0.019 0.060 0.060 0.19

279 TO(L)-TA(L) - 0.013 0.022 0.070 0.080 0.23

302 LA(L)+TA(L) or 2AII - 0.0033 0.009 0.040 0.040 0.10 354 TO(X),TO(L) - - - 0.44 0.56 1.10

366 TO(Γ) 1.00 1.00 1.00 1.00 1.00 1.00 376 LO(L) - - - 0.21 0.28 0.60 392 LO(X) - - - 0.35 0.48 0.60

403b LO(Γ) 0.39 0.52 0.54 0.54 0.63 0.66

420 2AIII - - - 0.11 0.12 0.30 431 2LA(L) - 0.042 0.053 0.12 0.14 0.38

453c TO(L)+TA(L) - 0.039 0.054 0.14 0.17 0.46

470 TO(X)+TA(X) - 0.036 0.058 0.18 0.22 0.76 550 GaN (E1(TO)) - - - - - 0.11 a. Relative to TO(Γ) mode at 365 cm-1 ; b. The specific frequencies for the observed LO(Γ) modes were 403, 402, 402, 402, 400 and 401.88 cm-1, respectively c. The specific frequencies for the observed TO(L)+TA(L) modes were 456, 455, 455, 453, and 452 cm-1, respectively

Figure III.6 Raman spectra of GaP nanowires film annealed at 750 °C in a gas stream of NH3 (g) and Ar (g) (1:5, v/v) with excitation at different laser wavelength. The red line describes the Raman spectrum obtained with 632.8 nm HeNe laser. The black line shows the Raman spectrum with 514.5 nm Ar+ laser as excitation source.

64 contrast to the data in Figure III.4b, the relative intensities of the TO(Γ) and LO(Γ) modes were not affected by changes in light polarization. In fact, the relative intensities of all observable spectral features for the N-alloyed nanowire in Figure III.4c were nominally invariant to the polarization conditions of the experiment.

III.3.4 Transmission Electron Microscopy Figure III.7 shows representative transmission electron micrographs of individual

GaP nanowires before and after annealing at T = 750 ºC in mixtures of NH3(g) and Ar(g). Figures III.7a and III.7f highlight the observed crystallinity in a GaP nanowire prior to exposure to NH3(g). The long axis of the GaP nanowires coincided with the [111] direction. Twinning along the entire length of the nanowires was observed, a common feature of zincblende semiconductor nanowires.26 The electron diffraction pattern was collected over the entire diameter of the nanowire along the [111] zone axis and indicated long range crystallinity. In order to investigate the affect of heating on crystallinity, GaP nanowires were annealed in pure Ar at 750 °C for 4h. Transmission electron micrographs of a representative nanowire treated in this manner are shown in Figure III.8. Annealing in Ar removes the sporadic twinning defects observed in GaP nanowires (as described in

Figure III.7a and III.7f). As compared to GaP nanowires treated with different NH3 dilutions (see Figure III.7), the Ar-treated GaP nanowires have core-shell structure instead of forming pocket domains. The well-patterned electron diffraction data shown in the inset of Figure III.8b suggests that the overall crystallinity of Ar-treated nanowires is preserved. Subsequent annealing at 750 ºC in pure Ar (g) effected removal of twinning defects but did not otherwise perturb the nanowire morphology. Figures III.7b and III.7g show the effect of annealing in dilute NH3(g) on crystallinity. GaP nanowires annealed in

1:20 v/v NH3(g):Ar(g) exhibited significant surface roughening, with disruptions along the sides of the nanowire. Both the overall uniform contrast in Figure III.7b and the well resolved electron diffraction pattern (taken along the <112> zone axis for Figures III.7g-j) in Figure III.7g indicated that the long-range crystalline zincblende order along the [111] direction was maintained for the majority of the GaP nanowire. The micrographs for GaP nanowires annealed with higher levels of NH3(g) showed the appearance of separate crystalline domains (lighter regions in Figures III.7c-e and III.7h-j). Quantitative detection of N in these regions was not possible due to the low sensitivity towards N of

a)b) c)d) e)

20 nm 20 nm 20 nm 20 nm 20 nm

f)g) h)i) j)

5nm 5nm 5nm 5nm 5nm

Figure III.7 Transmission electron micrographs of representative GaP nanowires after annealing at T = 750 ºC. (a, f) GaP nanowire without annealing; (b,g) NH3(g):Ar(g)(1:20 v/v); (c,h) NH3(g):Ar(g) (1:14 v/v); (d,i) NH3(g):Ar(g) (1:9 v/v); (e, j) NH3(g):Ar (g) (1:5 v/v).

Figure III.8 Transmission electron micrographs of a representative GaP nanowire annealed in pure argon at 750 °C for 4h: (a) low magnification and (b) high resolution micrographs, inset: selected area electron diffraction pattern.

Figure III.9 High resolution transmission electron micrograph of a representative GaP nanowire annealed in a gas stream of NH3 (g) and Ar (g) (1:1, v/v), two pararell lines with arrows describe local lattice orientation.

68 the Si (Li) detector used for X-ray energy dispersive. The collected electron diffraction patterns indicated that the majority of each nanowire retained a GaP zincblende lattice structure (without evidence of crystalline GaN or other crystal structures) and that the [111] axis remained parallel with the nanowire length. However, high resolution micrographs showed that the new domains were ordered with crystalline orientations different than the crystallinity of the rest of the nanowire (Figure III.9). At the highest concentrations of NH3 (g), annealing caused partial structural damage in the form of etched pits along the sides.

III.3.5 X-ray Photoelectron Spectroscopy X-ray photoelectron spectra were collected to assess the chemical enviorment of the GaP nanowire films after annealing in NH3(g). Figure III.10 presents representative high-resolution N 1s XP spectra for GaP nanowire films after various treatments as compared to untreated GaP nanowires and single crystalline GaN. The GaN sample exhibited a well resolved singlet at 397.3 eV indicative of N atoms bound to Ga atoms in 27 a tetrahedral coordination. In the absence of intentional treatment with NH3(g), the absence of spectral signatures between 392 and 402 eV for as-prepared GaP nanowires indicated a lack of detectable N species at the surface or within the near surface-region.

After annealing in NH3(g), the N1s spectra for GaP nanowires annealed in 1:9 v/v

NH3(g):Ar(g) showed a detectable and markedly increased level of N. The spectra for these GaP nanowire films showed a single, broad signal centered at the same binding energy of 397.3 eV as the lone peak in the spectrum for GaN. No separate peak was 28 observed at 400 eV or higher binding energies to suggest the presence of interstitial N2 29,30 or N species or oxidized NOx groups. The XP spectrum for a GaP nanowire film annealed in more dilute NH3(g) also showed only a single peak, centered at a slightly lower binding energy of 397.2 eV. In both spectra of GaP nanowire films exposed to

NH3(g), the N1s singlet was noticeably broader than the corresponding line width in the spectrum for GaN. Comparison of Ga 3d spectra (Figure III.11) for GaP(111)B wafers and as-prepared GaP nanowire films showed that the difference in scattering from planar and nanowire film interfaces caused a comparatively small line broadening (< 0.2 eV). The line broadening in Figure III.10 with respect to the N1s spectra for GaN and the

GaN

GaP1-xNx 1:1 (NH3:Ar) N(e)/ e GaP1-xNx 1:9 (NH3:Ar)

GaP

400 398 396 394

Binding Energy /eV

GaP Nanowires annealed at 750oC NH /Ar(1:5) 3

GaP Nanowires without annealing

N(e)/e GaP Wafer (111)B

24 22 20 18 Binding Energy (eV)

NH3(g)-treated GaP nanowire films was large (1.33 and 1.68 eV, respectively). Hence, the N1s spectra for GaP nanowires treated in dilute NH3(g) in Figure 10 indicated some heterogeneity in the local chemical environment of N atoms, but less than what has been 29-31 reported for N-enriched metal oxides (e.g. TiO2) after annealing in NH3(g).

III.3.6 Long Wavelength Photoelectrochemical Conversion The light-harvesting properties of the as-prepared GaP nanowire films were assessed in a ferrocene/ferrocenium non-aqueous regenerative photoelectrochemical cell. Previous measurements of n-GaP photoelectrodes showed that this electrolyte effects a large interfacial barrier height (>1.5 eV).32 Figure III.12a shows the wavelength dependence of the external quantum efficiency for photon to electrical energy conversion at short circuit for two types of n-GaP photoelectrodes. The lowest energy bandgap for GaP allows for absorption at λ ≤ 545 nm33. However, the planar n-GaP photoelectrode in Figure III.12a only exhibited optical conversion quantum yields exceeding 0.01 for λ < 460 nm, illustrative of capture/conversion efficiency losses incurred by low-aspect ratio photoelectrodes with short minority carrier diffusion lengths relative to the optical penetration depth of long wavelength excitation.34 Figure III.12a also contains data for n-GaP nanowire film photoelectrodes. In comparison, the nanowire films photoelectrodes showed superior photoresponse properties at long wavelengths, i.e. quantum yields ≥ 0.01 for λ < 545 nm. The enhanced photoresponses for nanowire films relative to planar photoelectrodes were consistent with the premise that the high-aspect ratio form factor decoupled light absorption and carrier collection and facilitated enhanced photon-to- charge conversion at energies near the bandgap.35,11 Figure III.12b summarizes the observed wavelength-dependent photoresponses for several types of n-GaP nanowire film electrodes. Each photoelectrode type was analyzed in triplicate, with the average values and standard deviations for each quantum yield measurement shown. In dry acetonitrile containing dissolved ferrocene/ferrocenium, as-prepared n-GaP nanowire film photoelectrodes showed quantum yields that were less than 0.01 for λ > 545 nm and progressively decreased at longer wavelengths. After annealing in NH3(g) at T = 750 ºC, the nanowire film electrodes exhibited different quantum yield wavelength dependencies at sub-bandgap wavelengths. Annealing in 1:20 v/v NH3(g):Ar(g) caused a systematic suppression of the quantum yield values in the

No Annealing 0.6 a) b) 0.020 0.015 NH3:Ar (1:20) NH :Ar (1:14) 0.5 0.03 0.010 3 = 600 nm

Yield 0.005 NH3:Ar (1:9)  0.000 m 0.4 at NH :Ar (1:5) Quantum Yield 0.00 0.05 0.10 0.15 0.20 3 0.02 Fraction of NH (g) 0.3 3 0.2 GaP NWs Electrode 0.01 0.1 GaP Planar Electrode External Quantu External Yield Quantum 0.0 0.00 350 400 450 500 550 600 650 700 550 600 650 700 Wavelength (nm) Wavelength (nm) Figure III.12 (a) Comparison of the wavelength dependence of the external quantum yields obtained with (open diamonds) a planar n-GaP and (closed circles) a n-GaP nanowire film photoelectrode immersed in a 1M LiClO4 acetonitrile solution containing 5 mM ferrocene (Fc) and trace amounts of ferrocenium (Fc+). The data were recorded at short circuit (i.e. E = 0 V vs E(Fc+/Fc)). (b) Wavelength dependence of the external quantum yields obtained with GaP nanowire films after annealing at T = 750°C with dilute NH3(g). This spectral range corresponds to wavelengths below the lowest bandgap energy of pure GaP. Inset: the external quantum yield at λ=600 nm plotted as a function of the concentration of NH3(g) used for the annealing step.

73 range of wavelengths shown in Figure III.12b. Annealing GaP nanowire films under concentrated NH3(g) conditions effected an increase in the overall quantum yield values.

Nanowire films exposed to an ambient containing 1:9 v/v NH3(g):Ar(g) during annealing showed the uniformly largest quantum yields between 550 nm and 680 nm. GaP nanowire films that had been annealed with 1:5 v/v NH3(g)/Ar(g) were noticeably less photoactive. The inset to Figure III.12b summarizes the observed trend for quantum yield values as a function of NH3(g) in the annealing ambient. The collective data indicated an optimum condition for annealing in the presence of NH3(g) for effecting increased light harvesting at longer wavelengths than the bandgap energy of pristine GaP.

III.4 Discussion The cumulative data support the contention that annealing GaP nanowires in effluent gas containing NH3(g) is a viable route for producing GaP1-xNx nanowires. The observation of reflections in the X-ray diffraction data for only zincblende GaP, the absence of GaN phonon modes in the Raman spectra of individual nanowires and nanowire films, and the selected area electron diffraction patterns with reflections for only zincblende GaP show that N is alloyed into the zincblende lattice of GaP without conversion to wurtzite GaN in dilute NH3 (g) at T = 750 °C. The specificity of XP spectra for surface analysis, in conjunction with the difficulties in quantitative analysis of non- 36 planar sample interfaces, preclude a precise determination of the GaP1-xNx alloy composition produced from each annealing trial. Nevertheless, high-resolution N1s X-ray photoelectron spectra do indicate that the near surface regions of the annealed GaP nanowires likely contain substitutional N atoms covalently coordinated rather than appreciable amounts of NOx oxides or interstitial N2 species. The transmission electron micrographs indicate that structural changes induced by annealing specifically in the presence of NH3(g) are inhomogeneous throughout the entire volume of the GaP nanowires. The transmission electron micrographs show that annealing GaP nanowires in effluent gas containing 5% v/v of NH3(g) largely disrupts just the surface. Annealing in effluent gas with higher contents of NH3(g) caused the emergence of distinct crystalline domains throughout the nanowires, consistent with the idea that these regions are rich with N atoms. The lack of sensitivity towards polarization conditions in the Raman spectra from individual GaP1-xNx nanowires indicates the new

74 spectral features induced by N incorporation do not have any specific orientation with respect to the long axis of the nanowire. The increase in the intensity of the new Raman spectral features tracks with the NH3(g) content in the effluent gas, as does the domain size evident in the transmission electron micrographs. Taken together, the transmission electron micrographs and the polarized Raman spectra thus suggest that N alloy formation is highly localized in regions randomly located throughout the nanowire. Many of the emergent features in the Raman spectra of GaP nanowires following high temperature treatment with NH3(g) have not previously been reported for GaP1-xNx alloys and therefore merit specific discussion. For GaP1-xNx alloys, several changes in the Raman spectra (with respect to the spectra for pristine GaP) have been previously noted, including the common observance of the LO(X) mode.14,37,38,39,40,41 Often referred to as the ‘X’ or defect mode, this spectral feature has been consistently observed in GaP1-xNx samples at room temperature irrespective of the method of sample preparation. Depending on the method of alloy preparation, a separate ‘local vibrational mode’ near 500 cm-1 which shows resonance enhancement has been noted37,38 and described as the vibrational frequency of isolated Ga-N bonds surrounded by Ga-P bonds.37 The observation of this mode typically requires a high concentration of N in the lattice and/or low temperatures (T < 20 K).39,41 The ‘local vibrational mode’ was not observed in the

GaP1-xNx nanowires at any excitation wavelength (λ = 514, 633, and 785 nm). A related -1 mode at 437 cm for N2 (rather than N) replacing P in the anionic sites in the lattice has 42 also been speculated for related GaAs1-xNx alloys but was also not observed in the

Raman spectra for the GaP1-xNx samples prepared here. However, the red shift of the LO(Γ) mode is consistent with a substitutional, rather than an interstitial, impurity in the 43 host lattice. Moreover, the room temperature Raman spectra of the NH3(g)-treated GaP nanowires did show numerous (up to 11) assignable Raman modes for GaP nanowires after annealing in NH3(g) that have not been observed previously for GaP1-xNx or related III-V N-alloys. The appearance of Raman signals at frequencies corresponding to phonon modes at the Brillouin zone boundaries suggests defect-activated phonon scattering from non-uniformly distributed impurities.18c Substitutions within the lattice destroy translational symmetry and relax the wave vector selection rule for Raman scattering (i.e. q ≠ 0), allowing the observation of formally forbidden acoustic and optical phonon

75 modes.16,17,18c Randomly substituted N atoms at anionic sites within GaP have been proposed previously to activate specifically phonon modes from the X point of the Brillouin zone,18c,38 consistent with the identification of TA(X), LO(X), and TO(X) modes (and their combinations) in the spectra reported here. The absence of a dependence on polarization of these features in the spectra in Figure III.4c stands in contrast to that 38 reported for epitaxial GaP1-xNx thin films but is consistent with a randomized incorporation within the identified domains in the transmission electron micrographs. The observation of several additional phonon modes from the L point of the Brillouin zone could indicate that multiple types of sites within the zincblende lattice are occupied by guest species or could simply mean that the defect-induced phonons arise from a point with a high density of states near both the X and L points.43 Additional (low temperature) Raman analyses are needed to better distinguish these possibilities. The available bulk and surface materials characterizations, in conjunction with separate photoelectrochemical data, indicate that all NH3(g)-based annealing conditions do not effect equivalent solar energy conversion properties in GaP. The optical properties inferred from the diffuse reflectance measurements indicate a steady increase in the absorbance of visible (sub-bandgap) light with an increasing fraction of NH3(g) in the effluent gas. These data are generally in accord with prior reports of slight bandgap 44 lowering (with respect to GaP) in GaP1-xNx alloys. However, the photoelectrochemical responses at these same sub-bandgap wavelengths did not show a similar monotonic increase in photoresponse at wavelengths longer than 545 nm. Instead, the photoelectrochemical data clearly indicate an optimal set of annealing conditions, with a steep decrease in quantum yields for optical energy conversion after annealing in gas mixtures with a high NH3(g) content. These observations imply that increased levels of N in GaP nanowires simultaneously increase the total absorbing power of light at λ > 545 nm but also decrease the quality of the electrical properties relevant to the capture of photogenerated carriers. For materials annealed with higher contents of NH3(g), the material degradation observed in the transmission electron micrographs and the preponderance of GaN-specific signatures in the Raman spectra and XRD patterns suggest that severe disruption of the GaP lattice negatively impacts the electrical properties. However, under processing conditions that facilitated uniform N incorporation

76 without major structural damage, the collected data do indicate that annealing in NH3(g) can increase the photoresponse of GaP nanowires at wavelengths below the formal bandgap energy. The findings of this work show that alloying N into GaP nanowires through annealing in dilute NH3(g) reproducibly red-shifts the onset of photocurrent by visible light by more than 100 nm. The possibility that additional post-synthetic routes for introducing N atoms into GaP nanostructures (e.g. annealing with hydrazine or at lower pressures) could result in similar or further enhancement in photoresponses at long wavelengths remains an open question. Since aspects including the electrically active dopant density,45 surface passivation, and nanowire packing density were not explicitly ‘optimized’ in these nanowire film photoelectrodes, the absolute energy conversion efficiencies of the photoelectrodes studied here were accordingly low in comparison to recent studies of nanostructured n-GaP which showed better cumulative energy conversion efficiencies.8 Nevertheless, the data shown here illustrate that optimally prepared GaP1-xNx nanowires could prove to be useful photoelectrode materials capable of sustaining large photovoltages and large photocurrents under normal solar insolation. Coupled with recent advances in methods to modify the surface chemistry and energetics of Ga-based III-V semiconductors through covalently attached surface groups,46 such materials should prove useful for photoelectrochemical applications for solar energy capture and conversion.

III.5 Conclusions GaP nanowire films were annealed in flowing gas containing varied levels of

NH3(g) at elevated temperatures, resulting in incorporation of N atoms into the lattice of the individual GaP nanowires. In effluent Ar(g) streams with a high (>20% v/v) NH3(g) content, X-ray diffraction patterns and Raman spectra separately indicated a conversion of zincblende GaP to wurtzite GaN at T ≥ 750 °C. At T = 750 °C, annealing in dilute

NH3(g) effected incorporation of substitutional N atoms into the GaP nanowires without compromising the zincblende structure. Changes in the polarized and non-polarized

Raman spectra following treatment with NH3(g) were consistent with disorder-activated phonon modes from substitutional N atoms at anionic sites in the zincblende lattice. Separate X-ray photoelectron spectroscopy further supported the notion that the

77 incorporated N is substitutional rather than interstitial or in the form of oxidizing NOx groups. The non uniformity did not appear to hamper sub-bandgap light absorption since increased absorptivity of visible light at wavelengths with energies less than the bandgap was observed in nanowire films treated with higher fractions of NH3(g). However, the quantum yields for optical to electrical energy conversion in a ferrocene-based regenerative photoelectrochemical cell did not show the same monotonic increase following annealing, implying that the electrical properties suffered after exposure to high levels of NH3(g). The present data thus identify a specific set of process conditions that improve the solar energy conversion properties at sub-bandgap wavelengths of GaP nanowire film photoelectrodes. Application of these findings on separately optimized GaP nanowire photoelectrodes should prove useful in assembling an efficient photoelectrochemical cell with the simultaneous capacity for appreciable photocurrents and photovoltages.

III.6 References 1. Bard, A. J.; Faulkner, L. R., Electrochemical Methods: Fundamentals and Applications. 2nd ed.; John Wiley & Sons, Inc.: 2001. 2. Heller, A.; Chang, K. C.; Miller, B., Photocurrent Spectroscopy of Semiconductor Electrodes in Liquid Junction Solar Cells. J. Am. Chem. Soc. 1978, 100, 684-688. 3. Fujishima, A.; Honda, K., Electrochemical Photolysis of Water at a Semiconductor Electrode. Nature 1972, 238, 37-38. 4. Turner, J. A., A Realizable Renewable Energy Future. Science 1999, 285, 687-689. 5. Bolton, J. R.; Strickler, S. J.; Connolly, J. S., Limiting and Realizable Efficiencies of Solar Photolysis of Water. Nature 1985, 316, 495-500. 6. Finklea, H. O., Semiconductor Electrodes. Elsevier: Amsterdam, 1988. 7. Aspnes, D. E.; Studna, A. A., Dielectric Functions and Optical Parameters of Si, Ge, GaP, GaAs, GaSb, InP, InAs, and InSb from 1.5 to 6.0 eV. Physical Reivew B 1983, 27, 985-1009. 8. Price, M. J.; Maldonado, S., Macroporous n-GaP in Nonaqueous Regerative Photoelectrochemical Cells. 2009, 113, 11988-11994. 9. Hagedorn, K.; Collins, S.; Maldonado, S., Preparation and Photoelectrochemical Activity of Macroporous p-GaP (100). J. Electrochem. Soc. 2010, 157, D588-D592. 10. Deutsch, T. G.; Koval, C. A.; Turner, J. A., III-V Nitride Epilayers for Photoelectrochemical Water Splitting: GaPN and GaAsPN. The Journal of Physical Chemistry B 2006, 110, 25297-25307. 11. MaioloIII, J. R.; Kayes, B. M.; Filler, M. A.; Putnam, M. C.; Kelzenberg, M. D.; Atwater, H. A.; Lewis, N. S., High Aspect Ratio Silicon Wire Array Photoelectrochemical Cells. J. Am. Chem. Soc. 2007, 129, 12346-12347.

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25. Wu, J.; Zhang, D.; Lu, Q.; Gutierrez, H. R.; Eklund, P. C., Polarized Raman Scattering from Single GaP Nanowires. Physical Reivew B 2010, 81, 165415 1-8. 26. (a) Johansson, J.; Karlsson, L. S.; Svensson, C. P. T.; Mårtensson, T.; Wacaser, B. A.; Deppert, K.; Samuelson, L.; Seifert, W., Structural Properties of <111> B-Oriented III-V Nanowires. Nat. Mater. 2006, 5, 574-580; (b) Algra, R. E.; Verheijen, M. A.; Feiner, L.- F.; Immink, G. G. W.; Theissmann, R.; Enckevort, W. J. P. v.; Vlieg, E.; Bakkers, E. P. A. M., Paired Twins and {112} Morphology in GaP Nanowires. Nano Lett. 2010, 10, 2349- 2356; (c) Verheijen, M. A.; Immink, G.; Smet, T. d.; Borgström, M. T.; Bakkers, E. P. A. M., Growth Kinetics of Heterostructured GaP-GaAs Nanowires. J. Am. Chem. Soc. 2006, 128, 1353-1359. 27. Maeda, K.; Teramura, K.; Saito, N.; Inoue, Y.; Domen, K., Photocatalytic Overall Water Splitting on Gallium Nitride Powder. Bull. Chem. Soc. Jpn. 2007, 80, 1004-1010. 28. Egelhoff(Jr.), W. F., N2 on Ni(100): Angular Dependence of the N1s XPS Peaks. Surf. Sci. 1984, 141, L324-328. 29. Mi, L.; Xu, P.; Wang, P. N., Experimental Study on the Bandgap Narrowings of TiO2 Films Calcined under N2 or NH3 Atmosphere Appl. Surf. Sci. 2008, 255, 2574-2580. 30. Cheung, S. H.; Nachimuthu, P.; Joly, A. G.; Engelhard, M. H.; Bowman, M. K.; Chambers, S. A., N Incorporation and Electronic Structure in N-doped TiO2 (110) Rutile. Surf. Sci. 2007, 601, 1754-1762. 31. Valentin, C. D.; Pacchioni, G.; Selloni, A.; Livraghi, S.; Giamello, E., Characterization of Paramagnetic Species in N-Doped TiO2 Powders by EPR Spectroscopy and DFT Calculations. The Journal of Physical Chemistry B 2005, 109, 11414-11419. 32. Mukherjee, J.; Erickson, B.; Maldonado, S., Physicochemical and Electrochemical Properties of Etched GaP (111)A and GaP (111)B Surfaces. J. Electrochem. Soc. 2010, 157, H487-H495. 33. Strehlow, W. H.; Cook, E. L., Compilation of Energy Band Gaps in Elemental and Binary Compound Semiconductors and Insulators. J. Phys. Chem. Ref. Data 1973, 2, 163-199. 34. Kayes, B. M.; Atwater, H. A.; Lewis, N. S., Comparison of the Device Physics Principles of Planar and Radial p-n Junction Nanorod Solar Cells. J. Appl. Phys. 2005, 97, 114302 1-11. 35. Spurgeon, J. M.; Atwater, H. A.; Lewis, N. S., A Comparison Between the Behavior of Nanorod Array and Planar Cd(Se,Te) Photoelectrodes. The Journal of Physical Chemistry C 2008, 112, 6186-6193. 36. Mohai, M.; Bertoti, I., Caculation of Overlayer Thickness on Curved Surfaces Based on XPS Intensities. Surf. Interface Anal. 2004, 36, 805-808. 37. Jackson, M. P.; Halsall, M. P.; Güngerich, M.; Klar, P. J.; Heimbrodt, W.; Geisz, J. F., Vibrational Properties of GaP and GaP1-xNx under hydrostatic Pressures up to 30 GPa. Physica Status Solidi (b) 2007, 244, 336-341. 38. Vorlíček, V.; Gregora, I.; Riede, V.; Neumann, H., Raman Scattering Study of GaP:N Epitaxial Layers. J. Phys. Chem. Solids 1988, 49, 797-805. 39. Kaneko, M.; Hashizume, T.; Odnoblyudov, V. A.; Tu, C. W., Electrical and Deep- Level Characterization of GaP1-xNx Grown by Gas-Source Molecular Beam Epitaxy. J. Appl. Phys. 2007, 101, 103707 1-5.

40. Pulzara-Mora, A.; Meléndez-Lira, M.; Jiménez-Sandoval, S.; Lopez-Lopez, M., Study of the Structural and Optical Properties of GaPN Thin Films Grown by Magnetron RF Sputtering. Vacuum 2006, 80, 468-474. 41. S.Yoon; Seong, M. J.; Geisz, J. F.; Duda, A.; Mascarenhas, A., Evolution of Electronc States in GaP1-xNx Studied by Resonant Raman Scattering Spectroscopy. Phys. Rev. B 2003, 67, 235209 1-4. 42. Ramsteiner, M.; Jiang, D. S.; Harris, J. S.; Ploog, K. H., Nonradiative Recombination Centers in Ga(As,N) and Their Annealing Behavior Studied by Raman Spectroscopy Applied Physics Letters 2004, 84, 1859-1861. 43. Cheong, H. M.; Zhang, Y.; Mascarenhas, A.; Geisz, J. F., Nitrogen-Induced Levels in GaAs1-xNx Studied with Resonant Raman Scattering. Phys. Rev. B 2000, 61, 13687- 13690. 44. (a) Bi, W. G.; Tu, C. W., N Incorporation in GaP and Band Gap Bowing of GaNxP1-x. Applied Physics Letters 1996, 69, 3710-3712; (b) Miyoshi, S.; Yaguchi, H.; Onabe, K.; Ito, R., Metalorganic Vapor Phase Epitaxy of GaP1-xNx Alloys on GaP. Applied Physics Letters 1993, 63, 3506-3608; (c) Baillargeon, J. N.; Cheng, K. Y.; Hofler, G. E.; Pearah, P. J.; Hsieh, K. C., Luminescence Quenching and the Formation of the GaP1-xN1-x Alloy in GaP with Increasing Nitrogen Content. Applied Physics Letters 1992, 60, 2540-2542. 45. Hagedorn, K.; Forgacs, C.; Collins, S.; Maldonado, S., Design Considerations for Nanowire Heterojunction in Solar Energy Conversion/Storage Applications. J. Phys. Chem. C 2010, 114, 12010-12017. 46. Mukherjee, J.; Peczonczyk, S.; Maldonado, S., Wet Chemical Functionalization of III-V Semiconductor Surfaces: Alkylation of Gallium Phosphide Using a Grignard Reaction Sequence. Langmuir 2010, 26, 10890-10896.

CHAPTER IV

Zinc Doping of Gallium Phosphide Nanowire Films

IV.1 Introduction This chapter reports on results from different synthetic attempts to dope GaP nanowire films. Specifically, several p-type doping methods are described and their impact on the resultant photoelectrochemical properties is shown. In Chapter II, the preparation of GaP nanowire films and their use as photoanodes in a photoelectrochemical cell was described. Since these nanowires are presumed to operate under depletion conditions, improvement of the quantum yield for photocurrent collection depends on the capacity of the nanowire to support a full depletion region. Specifically, the depletion region width, W, needs to be compact enough to fit within the nanowire radius. There are two approaches to achieve this condition. One strategy is to increase the diameters of the GaP nanowires. The other approach is to decrease the depletion region width through an increase in the dopant density as shown in Eq. IV.1. 1

2  W  0 b (IV.1) qN D where ε is the effective dielectric constant of the semiconductor, ε0 is the permittivity of free space, Φb is the equilibrium junction barrier height, q is the unsigned charge of an electron, ND is the doping density. The most commonly employed acceptor for p-type GaP is zinc (Zn). Zn is a shallow acceptor for III-V semiconductor and diffuses rapidly at high concentrations.2 In principle, introducing Zn into GaP nanowire films should render them p-type. If doped properly, the nanowires should be able to operate as photocathodes with a full depletion region at equilibrium. This work explores methods that either dope GaP nanowires while the nanowires are being formed or after their synthesis as described in Chapter II.

IV.2 Methods

IV.2.1 Zinc Doping Methods

IV.2.1.1 Zinc doping during GaP Nanowire Growth and Deposition In this method, Zn-containing source material is added directly with the GaP sublimation source powder. The CVD set-up is described in Chapter II. Three different zinc sources, including zinc powder (purified, J.T.Barker Inc.), zinc diphosphenediide

(ZnP2) (ball milled from zinc powder and phosphorus at a stoichiometric ratio of 1:2) and zinc phosphide (Zn3P2) (≥19% active phosphor (P) basis, Sigma-Aldrich) were used. GaP powder (99.99% metal basis, Sigma-Aldrich) was mixed thoroughly with Zn dopant source material until the color of the mixture was uniform. Table IV.1 summarizes the mass ratios used for the solid mixtures. The mixed powder was then placed on a quartz platform which was then inserted into a quartz tube (diameter 2.2 cm, length 66 cm) placed inside a single zone furnace. The growth substrate (a GaP(100) wafer with a thickness of 0.35 mm, EL-CAT.Inc) was placed in the cold zone area 7 inches away from the source. The system was pumped to 60 mTorr and then the furnace was heated to 800 oC for 30 min under static vacuum. Following, the system was cooled back to room temperature radiatively over 4 h.

Table IV.1 Mass ratio of zinc source and GaP Zinc Source : GaP Mass Ratio Zn : GaP 1:5, 1:9, 1:15 ZnP2 : GaP 1:1, 1: 5, 1:9, 1:11,1;20 Zn3P2 :GaP 85:62, 1:5,1:10

IV.2.1.2 P-type Doping of GaP Nanowires After Growth in a Closed End Tube

10 mg of ZnP2 were placed in a closed-end tube in the center of a single zone tube furnace. The GaP nanowire films were placed at a distance of 1 inch away and the tube was pumped to 60 mTorr for 20 min. The zinc source temperature ranged from 500 oC to 650 oC.

Figure IV.1 A schematic graph of the open tube system. In the top part of the scheme, the mass flow controller used to control the flow of incoming carrier gas is connected to one end of the tube. The other end of the tube is connected to the outlet of the gas. In the bottom part of the graph, the details of the placement of the zinc source and the GaP nanowire film are described.

Figure IV.2 A schematic graph of spin-on dopant procedure, including four steps: (1) coat wafer with spin-on glass; (2) bake the wafer to drive off solvents; (3) high temperature drive-in anneal; (4) remove glass layer with HF.

Table IV.2 Summary of Conditions for Zn Doping of a GaP Wafer in an Open Tube System. Zinc Mass Ar Flow Temperature Distance Dwell Time Source (mg) Rate (oC) ( inch) (min) (sccm) Zn 40.0 50 800 8.25 120 Zn 42.0 50 800 5.75 120 Zn(pellet) 40.0 50 800 5.75 120 Zn 45.0 50 800 5.75 120 Zn 45.0 100 800 5.75 90 Zn 25.0 100 800 5.75 90 Zn 45.0 100 800 3.75 90 Zn 45.0 150 700 7.00 90 ZnP2 25.0 100 800 5.75 90 ZnP2 27.3 100 800 5.75 90

Table IV.3 Summary of Conditions for Zn Doping of GaP Nanowire Films in an Open Tube System. Zinc Mass Ar Flow Rate Temperature Distance Dwell Time Source (mg) (sccm) (oC) ( inch) (min) Zn 43.0 50 750 5.75 120 Zn 40.0 50 700 5.75 120 Zn 41.5 50 750 5.75 120 Zn 40.0 150 700 7 120 Zn 40.0 150 700 5 120 Zn 40.0 150 700 4 120 ZnP2 20.0 150 650 5.75 60 ZnP2 20.0 150 700 5.75 60 ZnP2 20.0 150 650 5.75 60 ZnP2 20.0 150 650 5.75 60 ZnP2 20.0 150 700 5.75 60 ZnP2 20.0 150 700 5.75 90 ZnP2 16.2 20 700 1.0 60 ZnP2 15.0 20 700 1.0 60 ZnP2 19.5 20 700 1.0 60 ZnP2 20.1 20 700 1.0 90 ZnP2 20.0 20 750 1.0 60

IV.2.1.3 P-type Doping of GaP Nanowires After Growth in an Open End Tube The schematic graph of the open tube system is shown in Figure IV.1. The GaP nanowire film/GaP wafer was placed in the center zone of the furnace. The Zn source was placed upstream of the GaP nanowire films. Tables IV.2 and IV.3 summarize the conditions used for doping by this method.

IV.2.1.4 P-type Doping of GaP Nanowires with a Spin-on Glass Film Loaded with Dopants Silica films loaded with zinc phosphate were also used for doping. The procedure is described in Figure IV.2, where a glass film coats the material to be doped by casting a spin-on-glass solution. Different volumes of dopant glass, spin rate and times of spinning were used (Table IV.4). The glass films were annealed at different temperatures under Argon gas to drive in the dopant. Different anneal temperature and times were used. The final step is to remove the glass layer with an HF(aq) etch step. High concentration of HF (49% (aq) by weight) also etched GaP so 1% HF was used for removing glass layer.

IV.2.1.5 P-type Doping of GaP Nanowire Films through Thermal Drive In of Evaporated of Zn Films A layer of Zn metal was evaporated on one side of a wafer substrate as a Zn dopant source. This coated wafer was placed on the top of GaP nanowire film, roughly separate by 1 mm (Figure IV. 3). This construct was then placed in the center of the tube furnace and pumped down to 60 mTorr. GaP nanowire films were then heated to 600 oC, 650 oC and 700 oC under static vacuum to drive in Zn.

IV.2.1.6 P-type Doping of GaP Nanowire Films through Small Volume Annealing in an Ampoule 10 mg of Zn powder and a GaP nanowire film were placed in an ampoule. Before sealing, the ampoule was pump to 60 mTorr and refilled with nitrogen gas three times with a Schlenk line to a nitrogen pressure of 760 mmHg. Then the ampoule was placed in the center of a tube furnace and heated for a specific time (Table IV.4).

Table IV.4 Conditions for Zn Doping in an Ampoule.

mZn Annealing Temperature Annealing (mg) (oC) Time (min) 10 550 30 10 600 30 10 600 45 10 600 60

IV.2.2 Photoelectrochemical Measurements

IV.2.2.1 Electrode preparation Before use, 1500 Ȧ thickness of zinc was evaporated to the back side of a substrate for making an ohmic contact. After annealing at 600 oC for 60 min under 200 sccm Ar, the residue Zn was removed by an etch in concentrated H2SO4 (93-98%, doubly distilled, Sigma-Aldrich). These wafers were then used for nanowire growth. After the growth, the back side was scratched and etched with concentrated H2SO4 for 30 s, followed by a rinsing with distilled water. Electrodes were then prepared by attaching the back side of each substrate to a copper wire coil threaded through a glass tube with silver epoxy (GC Electronics) and sealed with inert epoxy (Hysol C).

IV.2.2.2 Current-voltage characterization Current-potential (J-E) characterization of doped GaP nanowire films was tested in aqueous solutions containing 20 mM methylviologen (1,1’-dimethyl-4,4’-bipyridinium dichloride, 98% , Sigma-Aldrich) in 1 M H2SO4 solution. A three-electrode configuration was utilized, with the nanowire films as a working electrodes, a platinum wire as a counter electrode, and a Ag/AgCl electrode as a reference. Linear sweep voltammetry with a CHI 760 potentiostat was used to record the photoelectrochemical responses. The scan rate was kept at 50 mV s-1.

IV.2.2.3 Spectral Response Measurement The same cell and apparatus as described in Chapters 2 and 3 was used for measurement of the wavelength-dependence of the quantum yield for photocurrent collection at short circuit.

]

Figure IV.6 (a) J-E characterization of GaP nanowires prepared by in-situ doping on degeneratively doped p-Si wafer with the mixture of Zn3P2 and GaP powder (1:10 by mass); (b) J-E characterization of GaP nanowires prepared by in-situ doping on degeneratively doped p-Si wafer with the mixture of ZnP2 and GaP powder (1:9 by mass).

IV.3 Results

IV.3.1 Zinc Doping

IV.3.1.1 P-type Doping of GaP Nanowire Films During Growth and Deposition in a Closed-End Tube When heating GaP, phosphorus tends to sublime easier than gallium and deposits on the tube. In order to compensate the loss of phosphorus, the zinc source including phosphorus is used for this purpose. In this section, two type of zinc phosphide are employed as zinc source: Zn3P2 and ZnP2. Figure IV.4a shows the top-down SEM images of GaP nanowires prepared when

Zn3P2 is mixed with the source GaP material. For a Zn3P2: GaP mass ratio of 1:10, the majority of the deposited material appears to have a high aspect ratio form factor. However, no spherical growth catalysts were apparent on the nanowire tips. Some nanowires were connected at their tips. Energy dispersive spectroscopic (EDS) measurements of the film in Figure IV.4b showed only signature for Ga, P and Si (the underlying substrate). However, no Zn signal was detected, suggesting a Zn concentration < 0.1% at%)3.

The results for other Zn3P2:GaP ratios are also shown. Figure IV.4c shows the

SEM images of GaP nanowires prepared from Zn3P2:GaP mixtures of 62:85 by mass.

When a higher Zn3P2 fraction was used, the morphology of resultant GaP nanowires was greatly impacted. Numerous bulky particles were formed along with scattered thin nanowires. EDS analysis of these films still showed no Zn signal. These results indicate that increasing the Zn3P2 did not directly translate in to GaP nanowires with higher Zn loading. Rather, the inclusion of Zn3P2 at these levels during growth significantly impacted the morphology of the resultant GaP nanowires.

At different Zn3P2:GaP ratios (shown in Table IV.1), only doping with a 1:10 by mass ratio maintained the nanowire morphology. J-E characterization of the prepared nanowire film with 1:10 Zn3P2: GaP ratio is shown in Figure IV.6a. The response appeared n-type, i.e. no anodic current was passed in the dark. When the film was illuminated by white light, there was no observable change in J-E response (red curve),

93 indicating the nanowire film prepared by this method is not clearly non-degenerately doped p-type.

In separate experiments, Zn3P2 was replaced with ZnP2 as the dopant source to be used during growth. Figure IV.5 shows SEM images of GaP nanowires synthesized with

1:9 ZnP2:GaP ratio by mass. The morphology of the obtained nanowires was similar to those obtained without adding ZnP2 in the starting material, indicating ZnP2 did not perturb the GaP VLS process. EDS measurements still showed no detectable Zn.

Increasing the ZnP2:GaP ratio to 1:5 by mass did affect the GaP VLS growth, resulting in clusters of nanowires. EDS analysis still showed that the major elements in this product were Ga, P and Si without any detectable zinc signal. From Table IV.1, the only ratio of

ZnP2:GaP that preserves the morphology of the desired GaP nanowires is 1:9 by mass. The J-E characterization of GaP nanowire films prepared with this ratio is shown in Figure IV.6b. As with Zn3P2, no significant photocurrent consistent with p-type doping was observed.

IV.3.1.2 P-type Doping of GaP Nanowires After Growth in a Large Volume Closed- Tube Under Vacuum Figure IV.7 shows the J-E characterization of GaP nanowire film after annealing in Zn vapor in a large volume closed end tube system. When the nanowire film was treated at 600 oC under vacuum for 20 min, the prepared film showed n-type behavior with no observable photoeffect (Figure IV.6b). Since Zn diffusion in GaP is thermally activated,4,5,6,7 increasing the temperature should facilitate more uniform incorporation on Zn,

 Ea kBT D  D0 e (IV.2) where D is the dopant diffusion coefficient, D0 is the measure of the frequency with which an atom attempts to make a jump over the barrier, Ea is the activation energy, and kB is the Boltzmann constant. When the nanowire films were treated at 650 oC under vacuum for 20 min, the GaP nanowire films still showed predominantly n-type character (Figure IV.7a) and no significant photocurrent density. This result suggests Zn doping by annealing in a large volume closed tube is an effective doping strategy for as-prepared GaP nanowires.

Figure IV.7 (a) J-E characterization of GaP nanowires on degenerately doped p-Si prepared by ex-situ doping in closed end tube system ( 9.6 mg ZnP2, GaP nanowire film 1inch away from source, P= 60 mTorr, T = 650 oC, t = 20 min); (b) J-E characterization of GaP nanowires on degenerately doped p-Si prepared by ex-situ doping in closed end tube system ( 9.6 mg ZnP2, GaP nanowire film 1inch away from source, P= 60 mTorr, T = 600 oC, t = 20 min).

IV.3.1.3 P-type Doping of GaP Nanowires After Growth in a Large Volume Open Tube at Ambient Pressure To ascertain the relevant process conditions, doping experiments were first performed on intrinsic GaP wafer samples. Figure IV.8 shows the J-E characterization of GaP wafers annealed in a Zn-rich atmosphere in an open end process tube. In contrast to the results listed above, all samples prepared in this manner showed obvious p-type J-E character, i.e. in the illumination condition, when the potential scans more negative than the open circuit potential, the cathodic current starts to increase until reaching the plateau. In the absence of illumination, as the potential scans more negative, no current was observed. When the potential goes more positive, the anodic current is observed. The effects of temperature, carrier gass (Ar) flow rate, position in the process tube, and annealing time were investigated. Table IV.6 listed the measurable photocurrent density from samples treated under various conditions. The data show that temperature and annealing time were strongly influential. For instance, by increasing the annealing temperature by 100 oC, the photocurrent density increased by nearly one order of magnitude from 0.012 mA cm-2 to 0.068 mA cm-2, implying a higher dopant density that better supports the prospects for GaP nanowires as photoelectrodes operating under depletion conditions. By further increasing the annealing time from 90 min to 120 min, the photocurrent density increased by a factor of five. The specific Zn source material was also influential. Replacing ZnP2 with pure Zn increased the photocurrent density to 0.068 mA cm-2. The flow rate of the carrier gas was also influential. Slower flow rates resulted in higher attainable photocurrent density. With these results for a wafer, analogous experiments were performed on GaP nanowire films. Annealing these GaP nanowires at 800 oC caused the color to turn from from yellow to purple, indicating a decomposition route not seen for the GaP wafers. Accordingly, the highest annealing temperature used with GaP nanowires was 750 oC in this study. J-E characterization of GaP nanowire films is shown in Figure IV 9. All three samples demonstrated identifiable electrochemical behaviors consistent with p-type character. Table IV.7 lists the photocurrent densities for nanowire samples treated und

Table IV.5 Results summary of Zn Doping of a GaP Wafer in an Open Tube System. Zinc Mass Ar Flow Temperature Distance Dwell Photocurrent Source (mg) Rate (oC) ( inch) Time (mA/cm2) (sccm) (min) Zn 40.0 50 800 8.25 120 0.28(-0.6V) Zn 42.0 50 800 5.75 120 0.23(-0.6V) Zn(pellet) 40.0 50 800 5.75 120 0.25(-0.8V) Zn 45.0 50 800 5.75 120 0.23(-0.8V) Zn 45.0 100 800 5.75 90 0.068 (-0.4V) Zn 25.0 100 800 5.75 90 0.046(-0.4V) Zn 45.0 100 800 3.75 90 0.083(-0.4V) Zn 45.0 150 700 7.00 90 0.012 (-0.6V) ZnP2 25.0 100 800 5.75 90 0.013(-0.4V) ZnP2 27.3 100 800 5.75 90 0.016(-0.4V)

Table IV.6 Results summary of Zn Doping of GaP Nanowire Films in an Open Tube System. Zinc Mass Ar Flow Temperature Distance Dwell Photocurrent Source (mg) Rate (oC) ( inch) Time density (sccm) (min) (mA/cm2) Zn 43.0 50 750 5.75 120 0.026(-0.2V) Zn 40.0 50 700 5.75 120 0.065(-0.4V) Zn 41.5 50 750 5.75 120 Not detectable Zn 40.0 150 700 7 120 0.0040(-0.4V) Zn 40.0 150 700 5 120 0.029(-0.4) Zn 40.0 150 700 4 120 0.030(-0.4) ZnP2 20.0 150 650 5.75 60 0.0031(-0.4V) ZnP2 20.0 150 700 5.75 60 0.0060(-0.8V) ZnP2 20.0 150 650 5.75 60 Not detectable ZnP2 20.0 150 650 5.75 60 Not detectable ZnP2 20.0 150 700 5.75 60 Not detectable ZnP2 16.2 20 700 1.0 60 0.169(-0.3V) ZnP2 15.0 20 700 1.0 60 Not detectable ZnP2 19.5 20 700 1.0 60 Not detectable ZnP2 20.1 20 700 1.0 90 0.085(-0.3V) ZnP2 20.0 20 750 1.0 60 0.086(-0.3V) ZnP2 20.0 150 700 5.75 90 Not detectable

0.2 a)0.2 b) ) ) 2 2

0.0 0.0

-0.2 dark -0.2 dark illuminated illuminated Current ( mA/cm Current Current ( mA/cm

-0.4 -0.4 0.6 0.4 0.2 0.0 -0.2 -0.4 -0.6 0.6 0.4 0.2 0.0 -0.2 -0.4 -0.6 Potential ( V ) Potential ( V )

0.2 c)0.2 d) ) ) 2 2

0.0 0.0 mA/cm  dark -0.2 -0.2 dark illuminated illuminated Current / Current Current / ( mA/cm -0.4 -0.4 0.6 0.4 0.2 0.0 -0.2 -0.4 -0.6 -0.8 0.6 0.4 0.2 0.0 -0.2 -0.4 -0.6 -0.8 Potential ( V ) Potential ( V )

Figure IV.8 J-E characterization of GaP wafer by ex-situ doping in open end tube system (a) 45.0 mg Zn, non-doped GaP wafer 7 inch away from source, T = 700 oC, t = 90 min, Ar flow: 150 sccm; (b) 25.0 mg ZnP2, non-doped GaP wafer 5.75 inches away from source, T = 800 oC, t = 90 min, Ar flow: 100 sccm; (c) 45.0 mg Zn, non-doped GaP wafer 5.75 inches away from source, T = 800 oC, t = 90 min, Ar flow: 100 sccm; (d) 45.0 mg Zn, non-doped GaP wafer 5.75 inches away from source, T = 800 oC, t = 120 min, Ar flow: 50 sccm.

0.008 0.4 )

2 a) b) 0.0 0.006 -0.4 -0.8 0.004 dark -1.2 illuminated -1.6 Current ( mA/cm 0.002 Quantum Yield -2.0 -2.4 0.000 0.6 0.4 0.2 0.0 -0.2 -0.4 350 400 450 500 550 Potential ( V ) Wavelength / nm

0.4 0.16 ) ) 2 2 0.0 c) 0.00 d)

-0.4 -0.16

-0.8 -0.32 dark dark illuminated illuminated -1.2 -0.48

-1.6 -0.64

CurrentmA/cm Density ( -2.0 CurrentmA/cm Density ( -0.80 0.4 0.2 0.0 -0.2 -0.4 0.4 0.2 0.0 -0.2 -0.4 Potential ( V ) Potential ( V )

Figure IV.9 J-E characterization of GaP nanowires on p-GaP prepared by ex-situ doping in open end tube system (a) 16.2 mg ZnP2, GaP nanowire film 1inch away from source, T = 700 oC, t = 60 min, Ar flow: 20 sccm ;(b) corresponding quantum yield measurement of GaP nanowires prepared by the condition in(a); (c) 20.1 mg ZnP2, GaP nanowire film o 1inch away from source, T = 700 C, t = 90 min, Ar flow: 20 sccm; (d) 20.0 mg ZnP2, GaP nanowire film 1inch away from source, T = 750 oC, t = 60 min, Ar flow: 20 sccm.

The effects of temperature, carrier gass (Ar) flow rate, position in the process tube, different conditions. The photocurrent density trend was not as obvious as for GaP wafers. Due to the variation in the starting nanowires (i.e. the kinking of nanowires varies from different growth ) in each batch, the optimal doping condition was not clear. The highest photocurrent density was achieved when the nanowire film was annealed at 700 oC for 60 min. At this condition the photocurrent density of the prepared film was 0.169 mA cm-2 at -0.3 V. Increasing the annealing time or increasing the annealing temperature decreased the photocurrent density considerably. For example, when the nanowire film was annealed at 700 oC for 90 min, the photocurrent density dropped to 0.085 mA cm-2 at -0.3 V. The effects of temperature, carrier gass (Ar) flow rate, position in the process tube, and annealing time were investigated. Table IV.6 listed the measurable photocurrent density from samples treated under various conditions. The data show that temperature and annealing time were strongly influential. For instance, by increasing the annealing temperature by 100 oC, the photocurrent density increased by nearly one order of magnitude from 0.012 mA cm-2 to 0.068 mA cm-2, implying a higher dopant density that better supports the prospects for GaP nanowires as photoelectrodes operating under depletion conditions. By further increasing the annealing time from 90 min to 120 min, the photocurrent density increased by a factor of five. The specific Zn source material was also influential. Replacing ZnP2 with pure Zn increased the photocurrent density to 0.068 mA cm-2. The flow rate of the carrier gas was also influential. Slower flow rates resulted in higher attainable photocurrent density. With these results for a wafer, analogous experiments were performed on GaP nanowire films. Annealing these GaP nanowires at 800 oC caused the color to turn from from yellow to purple, indicating a decomposition route not seen for the GaP wafers. Accordingly, the highest annealing temperature used with GaP nanowires was 750 oC in this study. J-E characterization of GaP nanowire films is shown in Figure IV 9. All three samples demonstrated identifiable electrochemical behaviors consistent with p-type character. Table IV.7 lists the photocurrent densities for nanowire samples treated under

100 different conditions. The photocurrent density trend was not as obvious as for GaP wafers. Due to the variation in the starting nanowires (i.e. the kinking of nanowires varies from different growth ) in each batch, the optimal doping condition was not clear. The highest photocurrent density was achieved when the nanowire film was annealed at 700 oC for 60 min. At this condition the photocurrent density of the prepared film was 0.169 mA cm-2 at -0.3 V. Increasing the annealing time or increasing the annealing temperature decreased the photocurrent density considerably. For example, when the nanowire film was annealed at 700 oC for 90 min, the photocurrent density dropped to 0.085 mA cm-2 at -0.3 V. All GaP nanowire samples presented in Figure IV.9 were grown on p-GaP wafer. The substrate wafer has non-negligible effects on the J-E characterization. Therefore the results just from the steady-state J-E response could be convoluted with contributions from the underlying substrate. To separate the contributions, an analysis of the spectral responsivity of the quantum yield for photocurrent collection was performed. The quantum yield values for planar GaP photoelectrodes is known to drop dramatically at wavelengths longer than the direct bandgap wavelength (440 nm), falling to essentially zero by the indirect bandgap wavelength (550 nm). As detailed previously, the loss in responsivity is primarily because the collection pathlengths for optical absorption are longer than the maximum collection distance (as set by diffusional transport and carrier recombination). For photoactive nanowires, the quantum yield can be high until the indirect bandgap (550 nm) because the collection distance is just the radius of the nanowire. The quantum yield measurement of the prepared sample with the largest photocurrent density suggests the photocurrent generated by long wavelength light was perhaps more from the underlying substrate than the nanowires (Figure .IV.9).

IV.3.1.4 P-type Doping of GaP Nanowires with a Spin-on Glass Film Loaded with Dopants The uniform coating of tall, thin GaP nanowires with a glass film is challenging but can be controlled by altering the times of spin-coating and the amount of spin-on- glass (SoG) used. Table IV.8 listed all the conditions explored here. Figure IV.10 shows the J-E characterization and corresponding quantum yield measurement of the GaP

101

Table IV.7 Results summary of different conditions of Spin on Glass procedures for GaP nanowires. Volume Number Spin Pre- Pre- Ar Drive Drive Ar Photo (μl of Speed Temp Time Flow in in Flow -current /time) Spin Casts (rpm) (oC) (min) Rate Temp Time Rate Density (sccm) (oC) (min) (sccm) (mA/cm2) 50 1 4000 200 5 100 750 120 100 Not detectable 20 1 4000 200 5 100 750 120 100 Not detectable 200 1 4000 200 5 100 750 120 100 Not detectable 200 1 4000 100 5 100 750 240 200 Not detectable 400 1 4000 100 5 100 750 240 200 Not detectable 600 1 4000 100 15 100 750 240 200 Not detectable 800 1 4000 100 15 100 750 240 200 Not detectable 500 1 4000 140 60 500 750 240 200 Not detectable 500 1 4000 140 60 500 750 120 200 Not detectable 500 1 4000 140 60 500 700 120 200 Not detectable 1000 1 4000 140 60 500 700 120 200 0.025 (-1.0V) 500 2 4000 140 60 500 700 240 200 0.0050 (-1.0V) 125 16 2000 140 60 500 700 120 200 0.0044 (-1.0V) 62.5 32 2000 140 60 500 700 120 200 0.0008 (-0.6V) 62.5 24 2000 140 60 500 700 120 200 0.0026 (-0.6V) 62.5 16 2000 140 60 500 700 120 200 0.0040 (-0.6V) 20 12 2000 140 60 500 700 120 200 0.0011 (-0.45V) 20 6 2000 140 60 500 700 120 200 Not detectable 20 4 2000 140 60 500 700 120 200 Not detectable

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0.006 a) 0.0025 0.004 b) ) 2 0.002 0.0020 0.000 0.0015 -0.002

-0.004 0.0010 dark -0.006 Yield Quantum 0.0005

Current ( mA/cm Current illuminated -0.008 0.0000 -0.010 0.2 0.0 -0.2 -0.4 -0.6 350 400 450 500 550 Potential ( V ) Wavelength / nm 0.006 0.0025 0.004 c) d) ) 2 0.002 0.0020

0.000 0.0015 -0.002 -0.004 0.0010 dark -0.006 Yield Quantum 0.0005

Current ( mA/cm illuminated -0.008 0.0000 -0.010 0.2 0.0 -0.2 -0.4 -0.6 350 400 450 500 550 Potential ( V ) Wavelength / nm

0.006 0.0025 0.004 e) f) ) 2 0.002 0.0020

0.000 0.0015 -0.002 0.0010 -0.004 dark

-0.006 illuminated Yield Quantum 0.0005 Current ( mA/cm -0.008 0.0000 -0.010 0.0 -0.2 -0.4 -0.6 350 400 450 500 550 Potential ( V ) Wavelength / nm

Figure IV.10 GaP nanowires on p-GaP were treated with different amount and times for spin coating, then were preannealed at 140 oC for 60 min with rate at 500 sccm Ar, and heated at 700 oC for 120 min with 200 sccm Ar. a) J-E characterization of GaP nanowires with 62.5 μl/time for 16 times; b) Quantum yield measurement of the sample corresponding to a), inset: SEM image of the prepared nanowires; c) J-E characterization of GaP nanowires with 62.5 μl/time for 24 times; d) Quantum yield measurement of the sample corresponding to c), inset: SEM image of the prepared nanowires; e) J-E characterization of GaP nanowires with 62.5μl/time for 32 times; f) Quantum yield measurement of the sample corresponding to e), inset: SEM image of the prepared nanowires.

103 nanowire film treated with different spin-coating conditions. Figure IV.10a shows the J-E characterization of GaP nanowire film spin coated with 16 casts of 62.5 μl of spin on glass. The J-E curve showed typical p-type character with photocurrent density around 0.0040 mA cm-2. Figure IV.10c shows the J-E curve of GaP nanowire film spin coated with 24 casts of 62.5 μl of spin on glass. The J-E response demonstrated p-type character with a photocurrent density around 0.0026 mA cm-2. Figure IV.10e shows the J-E characterization of GaP nanowire film spin coated with 32 casts of 62.5 μl of spin on glass. The nanowire film exhibited p-type character with a photocurrent density 0.0008 mA cm-2. With the same amount of spin on glass, increasing the number of casts decreased the photocurrent density of the nanowire film. Figures IV.10b, d, and f show the corresponding quantum yield measurements for the GaP nanowire films treated with different numbers of spin casts. Figure IV.10b shows that the largest quantum yield the film can achieve is 0.0025. The collective data indicate that increasing spin-coating steps leads to less photoactive nanowire films. Apparently, increasing the number SoG casts impacted the coverage of the spin on glass on the nanowires. The insets of Figure IV.10 are the cross- section SEM images of GaP nanowires with different cast numbers. The thickness of nanowire film changed from 15 μm to 5 μm with increasing spin-coating steps. It can been seen that most of the nanowires have their bottom part covered in the SoG, with only their top parts emerging outside of the SoG. For samples prepared with more spin-coating steps, the depth of the nanowires covered in SoG increased. Eventually all nanowires were completely buried in the SoG with sufficient spin-coating steps (32 times). At this condition the film became very dense and the thickness greatly decreased. Because of the high solubility of GaP in HF solution, the removal of the glass layer covering GaP nanowires has to be performed with low HF concentration (1%). SEM observation suggested at such low concentration the glass removal is not completed. Complete removal of the glass layer by HF etching without simultaneous etching of GaP was not achieved.

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Figure IV.11 (a) J-E characterization of GaP nanowire film on p-GaP with a silicon wafer (1 mm away) facing down treated at 650 oC for 2 h under vacuum. The Si wafer is spin coated with spin on glass 100 μl /time for 10 times, (b) Corresponding quantum yield measurement of the sample prepared by (a); (c) J-E characterization of GaP nanowire film annealed at 650 oC for 2h under vacuum with the Si wafer (1mm away) on the top. Si substrate was evaporated with a layer of zinc (100 nm), before annealing, GaP nanowire film was etched by gold etchant for 60 min; (d) Corresponding quantum yield measurement of the sample prepared by (c); (e) J-E characterization of GaP nanowire film prepared at the same annealing condition as (c). The difference is that the GaP nanowires were not etched; (f) Corresponding quantum yield measurement of the sample prepared by (e).

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IV.3.1.5 P-type Doping of GaP Nanowire Films through Sublimation of Zn from a Closely Spaced Spin on Glass Film In order to avoid the formation of a glass layer, attempts were made to separate the SoG from the GaP nanowires. SoG was spin-coated on one side of a Si substrate which was then placed 1 mm away from the top of GaP nanowire film by a glass spacer. Then the nanowire films were annealed at 650 oC for 2h under vacuum. Figure IV.11 shows the photoelectrochemical properties of the annealed GaP nanowire film treated under different conditions. All three samples showed p-type character. When the source was spin-coated on Si substrate, the photocurrent density was recorded approximately 0.01 mA cm-2(Figure IV.11a). The quantum yield measurement of this sample showed a peak around 450 nm with the highest value around 0.001, suggesting the recorded photoresponse originates from the nanowire instead of from the underlying GaP substrate. The photocurrent density dropped when a Zn metal film was used instead of the SoG. For this sample, the J-E characterization showed the largest photocurrent density achieved was around 0.004 mA cm-2. The corresponding quantum yield measurement showed a peak around the wavelength 500 nm, instead of around 450 nm. The J-E characterization showed that the achievable photocurrent density of this film was around 0.002 mA cm-2. The corresponding quantum yield measurement showed more complicated features. The first peak was around 375 nm with intensity around 0.00009 and the second peak was around 525 nm with intensity about 0.00005. The first feature resembles the profile of planar wafer. The second feature is more similar to the peak observed in (d), with a red shift of 15 nm. The J-E characterization of three doping conditions shows that the method with spin-coated spin on glass has the largest photocurrent density.

IV.3.1.6 P-Type Doping of GaP Nanowire films by Annealing in an Ampoule An ampoule was loaded with a GaP nanowire film sample and Zn powder. The ampoule was then pumped by vacuum pump, back filled with N2, and annealed. The photoelectrochemical properties of GaP nanowire film treated in this way at 600 oC for 60 min are shown in Figure IV.12. GaP nanowire films before and after this treatment are also shown. After the doping procedure, the color of GaP nanowire film changed from

106

Figure IV.12 (a) Optical image of a GaP nanowire film in an sealed ampoule with 0.020 g Zn powder under 1 atmosphere of N2; (b) Optical image of the GaP nanowire film of (a) annealed at 600 oC for 60 min; (c) J-E characterization of annealed GaP nanowire film; (d) Corresponding quantum yield measurement of GaP nanowire film of (c).

107 bright yellow to deep yellow. The J-E characterization (Figure IV.12c) showed p-type character with a photocurrent of 0.025 mA cm-2. The quantum yield measurement showed a peak near 450 nm with the highest value around 0.09. These results translated to the highest phoactivity of all explored doping methods. IV.4 Discussion

IV.4.1 P-type Doping During GaP Nanowire Growth and Deposition The samples obtained by in situ doping of GaP nanowires with a mixture of GaP and zinc source did not show any observable photocurrent with p-type response. It suggested that direct introducing zinc source into the starting materials did not successfully produce p-type GaP nanowires. Despite the unsuccessful doping results with this method, it is interesting to note that the morphology of the product is highly affected by the choice of starting materials. Increasing the Zn-source material content in the GaP sublimation source disrupted the GaP VLS growth process. A likely reason is that the Zn source sublimes before GaP, thereby altering the partial pressure of GaP during the VLS growth process.

IV.4.2 P-type Doping of GaP Nanowire Films by Annealing in a Large Volume, Closed End Tube At the annealing temperature > 600 oC under vacuum, the temperature gradient across the process tube caused a gradient in the Zn distribution during annealing. Presumably, Zn deposited quickly and early in the annealing at the cold ends of the tube, resulting in very short contact times for Zn to actually dope the GaP nanowires.

IV.4.1.3 P-type Doping of GaP Nanowire Films by Annealing in a Large Volume, Open End Tube Inert flowing gas has two benefits. First, by using high flow rate inert gas, the tube system was thoroughly purged, continuously maintaining a low O2 content in the tube. Second, at ambient pressure the sublimation temperature of all the Zn source materials was high, effectively lengthening the interaction time between zinc and the GaP nanowires. The low overall quantum yields measured for the doped GaP nanowires could arise for multiple reasons. 1)The large surface area of GaP nanowires translates in to

108 intolerably high surface recombination rates. 2) The underlying substrate introduces additional dopants which could cause a barrier for the carriers at the base of the nanowires. 3) The kinking of the nanowires lowers charge-carrier mobilities below the necessary value (~100 cm2 V-1 s-1). 4) The net dopant density could still be lower than what is necessary for the diameters of the GaP nanowires studied here.1

IV.4.1.4 Spin-on Dopant The challenge to use SoG for the doping of GaP nanowires is related to the different thermal expansion of GaP and SoG. When the glass is solidified, the glass tends to crack and contract, peeling from the substrate. This destroys the GaP nanowire film and exposes the underlying substrate.

IV.4.1.5 Doping with an Ampoule This method gives the highest quantum yield among these six methods considered in our study. The peak around 0.09 at 450 nm also suggests the photoresponse indeed originates from GaP nanowires. These results suggest this method is suitable for zinc doping, and deserves further optimization in future research. The synthesis with this method started with the purge of the ampoule at 1 atm pressure. Compared to other methods considered in our study, this initial pressure is quite high, which could greatly increase the sublime temperature of Zn. With enough Zn powder in the ampoule, the Zn vapor in the ampoule reaches the equilibrium partial pressure without any deposition on the ampoule wall. It provides continuous zinc rich vapor to GaP nanowires film during the whole annealing process. Therefore zinc doping is best achieved with this method.

IV.5 Conclusions Six doping methods have been described and discussed in this chapter. By evaluating these six methods, we have identified several key factors that affect the performance of the product. The pressure of zinc and the interaction time between the source and the substrate are identified to be critical to efficient doping. Morphology is also another important factor to be considered. The pressure and temperature during growth has a great effect on the quality of nanowires. Overall, it demonstrates effective p- type doping can be achieved by using the confined space doping method. To further

109 improve the quantum efficiency of nanowires, the quality of GaP nanowires needs to be improved by tuning the growth conditions.

IV.6 References 1. Hagedorn, K.; Forgacs, C.; Collins, S.; Maldonado, S., Design Considerations for Nanowire Heterojunction in Solar Energy Conversion/Storage Applications. The Journal of Physical Chemistry C 2010, 114 (27), 12010-12017. 2. Schubert, E. F., Doping in III-V Semiconductors. Cambridge University Press: Cambridge, 2005; p 606. 3. Kuisma-Kursula, P., Accuracy, precision and detection limits of SEM–WDS, SEM– EDS and PIXE in the multi-elemental analysis of medieval glass. X-Ray Spectrometry 2000, 29 (1), 111-118. 4. Sudlow, P. D.; Mottram, A.; Peaker, A. R., Doping gradients in layers of gallium phosphide grown by liquid epitaxy. J. Mater. Sci. 1972, 7 (2), 168-175. 5. Chang, L. L.; Pearson, G. L., Diffusion and Solubility of Zinc in Gallium Phosphide Single Crystals. J. Appl. Phys. 1964, 35 (2), 374-378. 6. Widmer, A. E.; Fehlmann, R., The diffusion of zinc in gallium phosphide under excess phosphorus pressure from a ZnP2 source. Solid-State Electron. 1971, 14 (5), 423-426. 7. Nygren, S. F.; Pearson, G. L., Zinc Diffusion into Gallium Phosphide under High and Low Phosphorus Overpressure. J. Electrochem. Soc. 1969, 116 (5), 648-654.

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CHAPTER V

Conclusions and Future Work

V.1 General Findings The research in this dissertation demonstrates three important points. (1) Gallium phosphide nanowires prepared by simple chemical vapor deposition exhibit appreciably increased photon-to-electron conversion efficiency in the visible wavelength region (450 nm to 550 nm). (2) The efficiency in the sub-bandgap region can be further improved by introducing nitrogen into GaP lattice to form GaP1-xNx. (3) Not all p-doping methods were equally effective for doping GaP nanowires. Improving the conversion efficiency, especially in the visible light region, was demonstrated in Chapter II. This result justifies the hypothesis that decoupling the light absorption and carrier collection directions can improve the carrier efficiency in a carrier- collection-limited material like GaP. The GaP nanowires directly prepared from the sublimation of raw GaP powder showed n-type character. By briefly etching with KI/I2 for 1 min, overall photoresponse characteristics were improved. Longer etching times did not necessarily increase the conversion efficiency, instead, with long etching time (in the scale of 15 min), the overall photoresponse deteriorated. Nitrogen alloyed GaP nanowires were successfully synthesized by annealing GaP nanowires in a flowing stream of NH3(g). The extent of nitrogen alloying was a strong function of the NH3/Ar ratio. The structural characterization illustrates nitrogen was incorporated non-uniformly throughout the nanowires at the anionic sites. The photoelectrochemical performance of nitrogen alloyed nanowires in non-aqueous regenerative cell demonstrated enhanced photoresponse in the sub-bandgap wavelength region (550-600 nm). The optimal nitrogen incorporation for best sub-bandgap enhancement was shown to be about 11%.

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To improve the overall efficiency of GaP nanowires, intentional doping was explored. For the purpose of making nanowire photocathodes, the as-prepared GaP nanowires were exposed to gaseous Zn. Although doping with Zn has been successfully demonstrated for planar GaP wafers, the controlled doping of nanowires was challenging. In Chapter IV, the results showed that the partial pressure of the gaseous Zn source and the total interaction time between the sample and Zn vapors were critical to successful p- doping. Overall, the present results suggest that annealing in a closed, small volume ampoule is the most reliable method.

V.2 Future Directions

V.2.1 P-type Doping of GaP Nanowires

V.2.1.1 Morphology control of GaP nanowires Preliminary data suggest the pressure and temperature are both critical to the growth of straight nanowires with a high quality. The disadvantage of current single heating zone furnace is that the temperature control of the substrate is convoluted with the source. And the pressure is under static vacuum and not accurately controlled. Therefore, for future work, it is essential to build a system to have better control over the temperature and pressure. One approach can be to incorporate a local heater heating the source while keeping the substrate in the center zone of the furnace. For pressure control, the inert gas flow can be introduced and a thrust valve can be helpful to maintain the pressure at certain value.

V.2.1.2 Single nanowire field-effect transistor measurement Chapter 4 focused on the preparing p-type zinc doped GaP nanowires. The characterization techniques for GaP nanowires in this thesis only involves bulk characterization methods, such as J-E responses and quantum yield measurement. These methods offer information on the average performance of an entire film array of nanowires. To quantitatively determine what specific doping conditions should be used for a given set of nanowires, detailed knowledge of the electrical properties of nanowires is essential. In this capacity, single nanowire field effect transistor1 (Figure V.1) can be

112

Source Drain a) Nanowire

Gate

Substrate

Oxide‐Insulator

Figure V.1 (a) A schematic sketch of the cross section of a single nanowire field effect transistor device; (b) A schematic sketch of a single nanowire field effect transistor device from top view.

113 used for this purpose. By monitoring the source to drain current (Isd) as a function of source to drain voltage (Vsd) at different gate bias (Vg), the type of doping, resistance and mobility data can be extracted. The doping density of the nanowire then can be calculated by the equation V.I, 1   (V.I) q(  n N d   p N a ) where ρ is the resistivity, μn and μp are the carrier mobility for electrons and holes respectively, Nd and Na are the doping density for electron and holes respectively.

V.2.2 Sensitized Solar Cell

V.2.2.1 Dye Sensitized Photoelectrochemistry GaP has a mid-sized bandgap and it absorbs light with shorter wavelength than 540 nm. Although the fraction of the absorption in solar irradiance is appreciable, there is still a portion of solar irradiation not being used. In chapter 3, the nitrogen alloying was one approach to extend the absorption of GaP more into the red end of the visible spectrum. Another strategy is to attach dyes that separately absorb sub-bandgap wavelength light. The valence band energetics of GaP favor sensitized hole injection due to its more negative valence bandage than the HOMO of most chromophores.(shown in Figure V.2a). This aspect facilitates sensitized hole injection (instead of electron injection) from a variety of common organic dyes. From a practical perspective, sensitized photocathodes in water are naturally suited for solar to chemical storage since the regeneration of the reduced dye can be coupled with reductive transformations in solution (shown in Figure V.3). The Maldonado group has recently shown sensitized hole injection into p-type GaP photoelectrodes is further facilitated when in operating under depletion conditions.2 For a planar photoelectrode, the dye loading is limited by low total available surface area. Semiconductor nanoparticles (<50 nm) provide ample surface areas but cannot typically support sufficient internal electric fields to help separate charges3. Instead, high aspect ratio nanowires with sufficient dopant density could be ideal platforms for the design of sensitized photocathodes.

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a b)

Figure V.2 (a) Dye-sensitized hole injection from a photoexcited chromophore at the surface of a phosphide semiconductor. (b) Dye-sensitized hole injection from a photoexcited chromophore at the surface of a phosphide semiconductor under depletion conditions. (source: Chitambar, M. et.al. “Dye-Sensitized Photocathodes: Efficient Light- Stimulated Hole Injection into p-GaP Under Depletion Conditions” J. Am. Chem. Soc. 2012, 134 (25), 10670-10681); (c) Schematic depiction of sensitized hole injection from a CdSe quantum dot into p-GaP.(source: Wang, Z. et.al. “Sensitization of p-GaP with CdSe Quantum Dots: Light-Stimulated Hole Injection. J. Am. Chem. Soc. 2013, 135 (25),9275- 9278).

115

CB Dye* e‐

hν VB H2O H2 + h Dye GaP

Figure V.3 A schematic graph of p-type dye sensitized photocathode for H2O reduction reaction.

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Sensitized nanowires/nanotubes in the conventional dye-sensitized photoanode design have been explored somewhat.4,5,6 They found an ordered topology increased the rate of electron transport and the nanowire electrode improved the quantum efficiency of DSCs. However, sensitized high aspect ratio p-GaP photoelectrodes have not yet been explored. This dissertation hints at routes towards making GaP nanowires that could be used in sensitized PEC systems. Refined doping control, in addition to surface modification strategies to bind dyes to GaP, are needed to realize such experiments. To bind dyes to GaP, there are several approaches. First, randomly physisorbed dyes on GaP are sufficient to show sensitization for study. Presumably, strongly chemisorbed dyes will prove more useful for an applied system. Previous work by the Maldonado group has shown that covalent Ga-C bonds can be formed at GaP surfaces through reaction with Grignard reagents.7,8 Similarly, P-O-C linkages seem promising for binding to GaP interfaces. It remains to be seen which route is the most fruitful or applicable to crystalline nanowire surfaces. This same surface modification will also be useful to determine exhaustively whether surface recombination was a predominant factor in the photoelectrochemical data shown in this thesis.

V.2.2.2 Quantum Dot Sensitized p-type GaP Quantum dots are attractive light harvesting materials.9 Different sizes of quantum dots can emit a series of different colors. Separate work from the Maldonado research group demonstrated CdSe quantum dots adsorbed on planar p-GaP photoelectrodes are a viable platform for studying sensitized hole injection. Performing these same experiments with high surface area p-GaP nanowire films will help increase the total quantum dot loading. The hope is that such a system may lead to photoelectrode responses with high internal and external quantum yields.

V.2.3 GaP Nanowires for PEC Cell P-type GaP planar wafer have been demonstrated to be a good photocathode for

H2. Although p-type GaP nanowire photocathode has been reported to be used for H2 reduction, the efficiencies to date are quite low10. For an efficient photocathode for hydrogen, an electrocatalyst is necessary. For example, adding a layer of platinum to a

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11,12 semiconductor photocathode greatly improved the efficacy for H2 generation. This same approach could be applied to the p-GaP nanowire films shown here. Stability in water is problematic for GaP. Adding a protective layer on the surface is therefore important. It has been demonstrated that the deposition of ultrathin oxide layers onto semiconductor surfaces can be one effective way.13 For instance, recent works show the deposition of TiO2 and MnO2 on n-type Si photoanodes for improving the stability.14,15 For the purpose of stability, the oxide does not need to be as conductive as conventional transparent conducting oxides due to the small dimension of the thin film.

V.3 References 1. Kim, D. M.; Jeong, Y.-H.; SpringerLink, Nanowire Field Effect Transistors: Principles and Applications. Springer New York : Imprint: Springer: New York, NY, 2014; p VI, 290 p. 196 illus., 98 illus. in color. 2. Chitambar, M.; Wang, Z.; Liu, Y.; Rockett, A.; Maldonado, S., Dye-Sensitized Photocathodes: Efficient Light-Stimulated Hole Injection into p-GaP Under Depletion Conditions. J. Am. Chem. Soc. 2012, 134 (25), 10670-10681. 3. Foley, J. M.; Price, M. J.; Feldblyum, J. I.; Maldonado, S., Analysis of the operation of thin nanowire photoelectrodes for solar energy conversion. Energy Environ. Sci. 2012, 5 (1), 5203-5220. 4. Law, M.; Greene, L. E.; Johnson, J. C.; Saykally, R.; Yang, P., Nanowire dye- sensitized solar cells. Nat Mater 2005, 4 (6), 455-459. 5. Baxter, J. B.; Aydil, E. S., Nanowire-based dye-sensitized solar cells. Applied Physics Letters 2005, 86 (5), -. 6. Tan, B.; Wu, Y., Dye-Sensitized Solar Cells Based on Anatase TiO2 Nanoparticle/Nanowire Composites. The Journal of Physical Chemistry B 2006, 110 (32), 15932-15938. 7. Mukherjee, J.; Peczonczyk, S.; Maldonado, S., Wet Chemical Functionalization of III- V Semiconductor Surfaces: Alkylation of Gallium Phosphide Using a Grignard Reaction Sequence. Langmuir 2010, 26, 10890-10896. 8. Peczonczyk, S. L.; Brown, E. S.; Maldonado, S., Secondary Functionalization of Allyl- Terminated GaP(111)A Surfaces via Heck Cross-Coupling Metathesis, Hydrosilylation, and Electrophilic Addition of Bromine. Langmuir 2013, 30 (1), 156-164. 9. Wang, Z.; Shakya, A.; Gu, J.; Lian, S.; Maldonado, S., Sensitization of p-GaP with CdSe Quantum Dots: Light-Stimulated Hole Injection. J. Am. Chem. Soc. 2013, 135 (25), 9275-9278. 10. Liu, C.; Sun, J. W.; Tang, J. Y.; Yang, P. D., Zn-Doped p-Type Gallium Phosphide Nanowire Photocathodes from a Surfactant-Free Solution Synthesis. Nano Lett. 2012, 12 (10), 5407-5411. 11. Peng, K.-Q.; Wang, X.; Wu, X.-L.; Lee, S.-T., Platinum Nanoparticle Decorated Silicon Nanowires for Efficient Solar Energy Conversion. Nano Lett. 2009, 9 (11), 3704- 3709.

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12. Dai, P.; Xie, J.; Mayer, M. T.; Yang, X.; Zhan, J.; Wang, D., Solar Hydrogen Generation by Silicon Nanowires Modified with Platinum Nanoparticle Catalysts by Atomic Layer Deposition. Angew. Chem. 2013, n/a-n/a. 13. McKone, J. R.; Lewis, N. S.; Gray, H. B., Will Solar-Driven Water-Splitting Devices See the Light of Day? Chem. Mater. 2013, 26 (1), 407-414. 14. Chen, Y. W.; Prange, J. D.; Dühnen, S.; Park, Y.; Gunji, M.; Chidsey, C. E. D.; McIntyre, P. C., Atomic layer-deposited tunnel oxide stabilizes silicon photoanodes for water oxidation. Nat Mater 2011, 10 (7), 539-544. 15. Strandwitz, N. C.; Comstock, D. J.; Grimm, R. L.; Nichols-Nielander, A. C.; Elam, J.; Lewis, N. S., Photoelectrochemical Behavior of n-type Si(100) Electrodes Coated with Thin Films of Manganese Oxide Grown by Atomic Layer Deposition. The Journal of Physical Chemistry C 2013, 117 (10), 4931-4936.

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