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Nanostructured Materials for Photovoltaics and Infrared Detection

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Title:
Nanostructured Materials for Photovoltaics and Infrared Detection
Creator:
Manders, Jesse R
Place of Publication:
[Gainesville, Fla.]
Florida
Publisher:
University of Florida
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Language:
english
Physical Description:
1 online resource (149 p.)

Thesis/Dissertation Information

Degree:
Doctorate ( Ph.D.)
Degree Grantor:
University of Florida
Degree Disciplines:
Materials Science and Engineering
Committee Chair:
SO,FRANKY FAT KEI
Committee Co-Chair:
PHILLPOT,SIMON R
Committee Members:
PEARTON,STEPHEN J
GILA,BRENT P
BOSMAN,GIJSBERTUS
Graduation Date:
5/3/2014

Subjects

Subjects / Keywords:
Electric potential ( jstor )
Electrons ( jstor )
Low temperature ( jstor )
Narrative devices ( jstor )
Ozone ( jstor )
Photometers ( jstor )
Photovoltaic cells ( jstor )
Polymers ( jstor )
Quantum dots ( jstor )
Solar spectra ( jstor )
Materials Science and Engineering -- Dissertations, Academic -- UF
energy -- infrared -- photodetectors -- photovoltaics -- polymers -- qds -- semiconductors
Genre:
bibliography ( marcgt )
theses ( marcgt )
government publication (state, provincial, terriorial, dependent) ( marcgt )
born-digital ( sobekcm )
Electronic Thesis or Dissertation
Materials Science and Engineering thesis, Ph.D.

Notes

Abstract:
Organic semiconductors and inorganic quantum dots are becoming topics of intense research due to their promising optical and electronic properties as well as their amenability to large-scale high-throughput solution-based processing methods. Organic solar cells are rapidly improving in performance with the promise of lightweight, functional and aesthetically pleasing applications. Inorganic quantum dots are especially interesting materials for next-generation optoelectronic sensors, particularly in the infrared spectrum, due to their easily tunable bandgap and controllable transport properties. First, we developed a new understanding of the interfacial properties of solution-processed NiO hole transport layers (HTLs) in polymer solar cells. Polymer solar cells were fabricated with NiO HTLs with a power conversion efficiency (PCE) of 7.8%, a 14% improvement over the reference devices. The improvement is due to an optical resonance shift in the solar cells, and an improvement in the morphology of the photoactive layer near the HTL/active layer interface that led to increased charge extraction over the reference devices. Additionally, the solar cells with NiO were more stable in air than the control devices. In the next section, we developed a new low-temperature fabrication process to incorporate solution-processed NiO HTLs into polymer solar cells. By pre-heating the NiO precursor films to a low temperature, then oxidizing the films, we formed NiO at temperatures low enough to be compatible with plastic substrates envisioned for roll-to-roll (R2R) processing. Solar cells fabricated on standard glass substrates showed comparable performance to the community standard devices and were slightly more stable in air. As a proof of concept, solar cells were fabricated on plastic substrates using this new method of fabricating NiO. The cells showed favorable performance, comparable with the community standard polymer solar cells typically fabricated on glass. In the final section, we demonstrated, for the first time, an all solution-processed, all-inorganic multispectral photodetector. The devices utilized solution-derived NiO and colloidal ZnO nanoparticles as charge blocking layers, and PbS quantum dots as the photoactive layer. Due to the wide bandgaps and favorable charge blocking ability of the oxides, the photodetectors produced extremely low dark current and noise levels. They also showed large linear dynamic range, adequate bandwidth for imaging applications, and a specific detectivity comparable to, or better than, some commercially available photodetectors. Additionally, because of the double heterojunction formed by oxides, the devices were very stable in air, remaining unchanged for five months. ( en )
General Note:
In the series University of Florida Digital Collections.
General Note:
Includes vita.
Bibliography:
Includes bibliographical references.
Source of Description:
Description based on online resource; title from PDF title page.
Source of Description:
This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Thesis:
Thesis (Ph.D.)--University of Florida, 2014.
Local:
Adviser: SO,FRANKY FAT KEI.
Local:
Co-adviser: PHILLPOT,SIMON R.
Electronic Access:
RESTRICTED TO UF STUDENTS, STAFF, FACULTY, AND ON-CAMPUS USE UNTIL 2014-11-30
Statement of Responsibility:
by Jesse R Manders.

Record Information

Source Institution:
UFRGP
Rights Management:
Applicable rights reserved.
Embargo Date:
11/30/2014
Resource Identifier:
907294884 ( OCLC )
Classification:
LD1780 2014 ( lcc )

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NANOSTRUCTURED MATERIALS FOR PHOTOVOLTAICS AND INFRARED DETECTION By JESSE ROBERT MANDERS A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2014

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2014 Jesse Robert Manders

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To my family and friends

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4 ACKNOWLEDGMENTS First and foremost, I give thanks to my advisor, Dr. Franky So. Without his generous support and guidance, I would not have had a successful career at the University of Florida. I have not only learned beautiful science and engi navigate the administrative and political waters of working in our field. Also, thanks go to my dissertation committee members Dr. Gijs Bosman, with whose lab I successfully collaborated, Dr. Stephen Pearton, Dr. Brent Gila, and Dr. Simon Phillpot, who accompanied me on a life changing trip to London for the World Lecture Competition in 2012. Thanks go to all of my labmates for supporting me and always being willing to h elp me omplete my task. Each of you have influenced me in your own way, and I am grateful for that. First, I thank Dr. Mike Hartel, the first member of the group I met when I came on my visit weekend in 2009, for showing me that grad uate students can be passionat e about their work and have fu n outside the lab, and encouraging me to come to Florida. It was one of the best decisions I have made in my life and I have you to thank. I especially thank Dr. Kaushik Roy Choudhury and Dr. Galileo Sarasqueta for introducing me to quantum dots and infrared technologies. I am eternally grateful to Dr. Sai Chen, who reenergized my career and taught me how to be a better scientist. Special thanks go to Erik Klump, Danae Constantinou, and Tzung H an Lai, my conference travel buddy, for having fun and being productive with me at the same time. without several teachers and professors in high school and college. I thank Mrs. Akridge and Mrs. Buddendeck at Centerville High School for inspiring me and affirming my love for chemistry and science in general. Thanks go to my physics and math professors at my alma mater, Miami University for spending countless hours with

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5 To my fraternity advisor and friend Professor Jerry Miller, my deepest gratitude is due for always believing in me as a student, a man and a leader. Your guidance and wisdom h as influenced me beyond words. I would not have had a successful career in the lab without the love and support of my friends I made outside the lab. From my neighbors and friends from Hillel and Chabad, to my fellow Miami University (Ohio) transplants and favorite cohort in the Student Personnel in Higher Education program, I will never forget the great times we had together and know we will continue to make great new memories. I thank my roommates Bryan and Lee for all the great times we shared, especiall y Roommate Wednesday and family dinners I also owe my teammates on the UF Club Vo lleyball team and my intramural volleyball champion teams thanks for giving me a fun and competitive outlet so I could stay young And of course I am grateful to my dear fri ends and co instructors at the Gator Sals a Club and Ritmo y Sabor Salsa for infusing Latin culture in my life and igniting a lifelong passion. I would not have come this f ar in my life without my mom, dad and brother. You have been my inspiration and guidi ng light my whole life. You believed in me when I struggled, and celebrated with me when I succeeded. You are the one constant in my ever changing life. I am humbled by your uncompromising love, dedication and support E ven though the content of this thesi s is very different from that w hich you have made your careers you have never lost interest in my success and happiness I will work my hardest to continue to make you proud personally and professionally, for the rest of my life

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6 TABLE OF CONTENTS page ACKNOWLEDGMENTS ................................ ................................ ................................ ............... 4 LIST OF TABLES ................................ ................................ ................................ ........................... 9 LIST OF FIGURES ................................ ................................ ................................ ....................... 10 ABSTRACT ................................ ................................ ................................ ................................ ... 17 CHAPTER 1 INTRODUCTION ................................ ................................ ................................ .................. 19 1.1 Organic Semiconductors ................................ ................................ ................................ ... 19 1. 1.1 What Are Organic Semiconductors? ................................ ................................ ...... 19 1.1.2 Electronic Structure of Organic Semiconductors ................................ ................... 21 1.1.2.1 Molecular orbital construction ................................ ................................ ..... 21 1.1.2.2 Excitons ................................ ................................ ................................ ........ 22 1.1.2.3 Charge transport ................................ ................................ ........................... 24 1.2 Organic Photovoltaics ................................ ................................ ................................ ....... 25 1.2.1 Background on Orga nic Photovoltaics ................................ ................................ ... 25 1.2.2 Measurement and Fundamental Electrical Characteristics ................................ ..... 26 1.2.3 Organic Solar Cell Device Structures ................................ ................................ ..... 28 1.2.4 The Importance of Interfaces ................................ ................................ .................. 30 1.3 Inorganic Nanocrystals and Quantum Dots ................................ ................................ ...... 31 1.3.1 Fundamentals of Nanocrystals and Quantum Dots ................................ ................ 31 1.3.2 Synthesis of Quantum Dots and Nanocrystals ................................ ....................... 32 1.3.3 Growth Mechanism of Quantum Dots and Nanocrystals ................................ ....... 34 1.3.4 Optical and Electronic Properties ................................ ................................ ........... 36 1.3.5 Charge Transport and Excitons in Quantum Dots ................................ .................. 38 1.4 Quantum Dot Photodetectors ................................ ................................ ............................ 41 1.4.1 Fundamentals of Photodetectors ................................ ................................ ............. 41 1.4.2 Performance Characteristics ................................ ................................ ................... 43 1.4.3 Quantum Dot Photodetector Device Structures ................................ ..................... 45 1.5 Figures ................................ ................................ ................................ .............................. 47 2 SOLUTION PROCE SSED NICKEL OXIDE HOLE TRANSPORT LAYERS IN HIGH EFFICIENCY PHOTOVOLTAIC CELLS ................................ ................................ 57 2.1 Background ................................ ................................ ................................ ....................... 57 2.2 Results and Discussion ................................ ................................ ................................ ..... 58 2.2.1 NiO Precursor Composition ................................ ................................ ................... 58 2.2.2 Characterization of NiO Films ................................ ................................ ............... 60 2.2.2.1 Optical properties ................................ ................................ ......................... 60 2.2.2.2 Electronic properties ................................ ................................ .................... 61

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7 2.2.3 Photovoltaic Cells with Solution Processed Nickel Oxide ................................ .... 62 2.2.3.1 Solar cell characteristics ................................ ................................ ............... 63 2.2.3.2 Analysis of solar cell performance ................................ ............................... 64 2.2.3.3 Stability of solar cells ................................ ................................ ................... 69 2.3 Conclusions ................................ ................................ ................................ ....................... 69 2.4 Experime ntal Procedure ................................ ................................ ................................ .... 70 2.5 Figures and Tables ................................ ................................ ................................ ............ 72 3 NOVEL LOW TEMPERATURE ROUTE TO SOLUTION PROCESSED NICKEL OXIDE HOLE TRANSPORT LAYERS IN POLYMER PHOTOVOLTAICS .................... 83 3.1 Background ................................ ................................ ................................ ....................... 83 3.2 Results an d Discussion ................................ ................................ ................................ ..... 86 3.2.1 Solar Cells from Low Temperature NiO Fabrication Part 1 ............................... 86 3.2.2 Low Temperature NiO Precursor ................................ ................................ ........... 87 3.2.3 Solar Cells from Low Temperature NiO Fabrication Part 2 ............................... 89 3.2.3.1 Device performance ................................ ................................ ..................... 89 3.2.3.2 Elucidating the UV O 3 treatment process ................................ .................... 89 3.2.3.3 Comparing low temperature NiO solar cells to standards ............................ 92 3.2.3.4 Air stability of the solar cells ................................ ................................ ....... 93 3.2.4 Solar Cells Fabricated on Plastic Substrates ................................ .......................... 93 3.3 Conclusions ................................ ................................ ................................ ....................... 94 3.4 Experimental Procedure ................................ ................................ ................................ .... 94 3.5 Figures and Tables ................................ ................................ ................................ ............ 97 4 AIR STABLE MULTISPECTRAL PHOTODETECTORS WITH LOW NOISE MADE FROM ALL SOLUTION PROCESSED INORGANIC SEMICONDUCTORS ................ 110 4.1 Background ................................ ................................ ................................ ..................... 110 4.2 Materials Properties ................................ ................................ ................................ ........ 112 4.2.1 PbS Quantum Dots ................................ ................................ ............................... 112 4.2. 2 Zinc Oxide Nanoparticles ................................ ................................ ..................... 113 4.2.3 Nickel Oxide Films ................................ ................................ ............................... 113 4.3 Photodetector Performance ................................ ................................ ............................. 114 4.3.1 External Quantum Efficiency and Responsivity ................................ .................. 1 14 4.3.2 Speed and Bandwidth ................................ ................................ ........................... 114 4.3.3 Noise Levels and Detectivity ................................ ................................ ................ 115 4.3.4 Linear Dynamic Range ................................ ................................ ......................... 119 4.3.5 Air Stability ................................ ................................ ................................ .......... 119 4.4 Conc lusions ................................ ................................ ................................ ..................... 120 4.5 Experimental Procedure ................................ ................................ ................................ .. 120 4.6 Figures ................................ ................................ ................................ ............................ 123 5 CONCLUSIONS AND FUTURE WORK ................................ ................................ ........... 133 5.1 Summary ................................ ................................ ................................ ......................... 133 5.2 Outlook and Future Studies ................................ ................................ ............................ 134

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8 APPENDIX A ORGANIC MOLECULAR ST RUCTURES ................................ ................................ ........ 137 B LIST OF PATENTS, PUBLICATIONS, AND PRESENTATIONS ................................ ... 138 Patents ................................ ................................ ................................ ................................ ... 138 Peer Reviewed Publications ................................ ................................ ................................ 138 Conference Presentations ................................ ................................ ................................ ...... 139 LI ST OF REFERENCES ................................ ................................ ................................ ............. 140 BIOGRAPHICAL SKETCH ................................ ................................ ................................ ....... 148

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9 LIST OF TABLES Table page 2 1 Performance of solar cells fabricated at various NiO processing temperatures. One standard deviation is reported in parenthesis. ................................ ................................ .... 82 2 2 Device characteristics of the solar cells fabricated in this study. One standard deviation is reported in parenthesis. S eries resistance was calculated using the illuminated J V values converging to V oc Shunt resistance was calculated using illuminated J V values converging to J sc ................................ ................................ .......... 82 3 1 Performance of solar cells fabricated with PEDOT:PSS, high temperature processed NiO, and the low temperature processed NiO from Precursor B. One standard deviation is reported in parenthesis. ................................ ................................ ................. 109

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10 LIST OF FIGURES Figure page 1 1 Three classes of organic molecules. ................................ ................................ ................... 47 1 2 Diagrams showing sp 2 hybridization with and bonding, and the energy level diagram for sp 2 hybridization ................................ ................................ ............................ 47 1 3 A diagram showing sp 2 hybridization and conjugation in a benzene ring. Image extracted from wikicommons. ................................ ................................ ........................... 48 1 4 A diagram showing the three types of excitons in solid materials and their energetic distribution in the bandgap. Images extracted and modified from MIT Open Courseware lec tures on Organic Electronics. ................................ ................................ .... 48 1 5 An electron/polaron hopping between potential wells. ................................ ...................... 49 1 6 The AM 1.5G solar spectrum. ................................ ................................ ............................ 49 1 7 A generic current density voltage curve in the dark and under photoexcitation showing the short circuit current, open circuit voltage, and maximum power p oint. ........ 50 1 8 Device structures for the bilayer planar heterojunction solar cell and the modern bulk heterojunction solar ce ll. ................................ ................................ ................................ .... 50 1 9 Synthesis setup for colloidal quantum dots. A three necked flask is loaded with one precursor solution under Ar flow w ith a thermocouple to track the temperature. The second precursor is injected, then extracted with a syringe ................................ .............. 51 1 10 The chemical processes involved in the nucleation and growth of colloidal quantum dots in the hot colloidal injection technique ................................ ................................ ..... 52 1 11 Free energy considerations in the nucleation phase of colloidal quantum dot synthesis. The plot shows the interplay between the surface free energy cost and volume free energy gain by forming nuclei ................................ ................................ ...... 52 1 12 The relationship between growth rate and size for different precursor solution concentrations. The critical size changes with precursor solution concentration ............. 53 1 13 The energy level evolution as atoms are assembled from one single cation anion pair to the bulk band structure. The band structure of quantum dots lies b etween the molecular orbitals and bulk band structure. ................................ ................................ ....... 53 1 14 Tunability of the quantum dot bandgap is shown in different batches of PbSe synthesized for quantum dot based light emitting diodes. The absorption edge changes as a function of quantum dot size ................................ ................................ ........ 54

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11 1 15 A schematic of an electron hopping between discreet energy states. This simulates the hopping transport mechanism active in quantum dot transport. ................................ .. 54 1 16 This plot shows the working wavelengths for all the common infrared photodetector materials. MCT = HgCdTe. Figure extracted from Teledyne Judson Technologies website ................................ ................................ ................................ .............................. 55 1 17 A schematic of the complex structure of expitaxially grown infrared detectors in a focal plane array mounted to the silicon read out circuit by indium bump bonding. The Si ROIC is typically a CMOS circuit, involving many processing steps by itself. .... 55 1 18 A comparison of the lateral channel struc ture and vertical stack device structure for photodetectors. The structure on the left is the lateral channel structure common in photoconductors. The structure on the right is the vertical stack structure common in photodiodes without gain. ................................ ................................ ................................ .. 56 2 1 High resolution electrospray ionization mass spectrum of the precursor compound. The peak family shown corresponds to the isotopic distribution of nickel; the ion at m/z = 239.0520 contains the most common isotope, Ni 58. The inset is the proposed structure of the ion. ................................ ................................ ................................ ............ 72 2 2 Theoretical high resolution mass spectrum for [Ni(MEA) 2 (OAc)] + This spectrum is in agreement with the experimentally acquired pattern for the [Ni(MEA) 2 (OAc)] + ion, confirming the unique mass assignment. ................................ ................................ .... 72 2 3 Absorption spectra of the transformation of the precursor into NiO. The material transforms from the precursor to a full nickel oxide film by thermal decomposition of the precursor materials, as is shown by the elimination of the visible and IR absorption peaks, and the appearance of the UV band edge of NiO. ................................ 73 2 4 Transmission spectrum of 5 nm thin NiO films is compared to 30 nm thin PEDOT:PSS films used in the optimized solar cells. The thin NiO films are more transparent than the PEDOT:PSS films used in solar cells at waveleng ths longer than 590 nm. ................................ ................................ ................................ .............................. 73 2 5 High resolution Ni 2p 3/2 XPS acquisition for NiO. The spectrum shows four contributions one from Ni 2 + in the octahedral NiO configuration at low binding energy (red), one from hydroxylated or defective NiO at an intermediate binding energy (blue), one from nickel oxyhydroxide (blueish green), and a high energy peak from a shake up process (light gree n) in the NiO lattice at the highest binding energy. ................................ ................................ ................................ ............................... 74 2 6 High resolution O 1s XPS acquisition for NiO. The spectrum shows three contributions one from O 2 in the octahedral NiO configuration at low binding energy (red) and one from hydroxylated or defective NiO (blue) and one from an oxyhydroxide at a higher binding energy (blueish green). ................................ ............... 74

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12 2 7 GIXRD spectrum of solution derived NiO, heated to 275 o C for 45 minutes on a glass substrate. The bulk stoichiometric NiO diffraction pattern is also show n for comparison. The shift of the experimentally acquired GIXRD spectrum is indicative of crystallographic vacancies and interstitials. ................................ ................................ .. 75 2 8 Angle resolved XPS spectra for the Ni 2p3/2 and O 1s states. At low take off angles, the peaks corresponding to hydroxylated (blue) and oxyhydroxylated (bright green) NiO increase in intensity relative to the NiO peak (dark red), correspondin g to extreme surface positions for the hydroxylation ................................ ............................... 75 2 9 Device structure and energy levels with respect to vacuum of the materials used in the photovoltaic cells. ................................ ................................ ................................ ........ 76 2 10 Illuminated current density voltage ( J V ) characteristics of the solar cells with NiO fabricated at different temperatures. After the full formation of NiO at 275 o C the solar cells perform optimally, with a PCE of 7.8%. ................................ .......................... 76 2 11 J V characteristics of solar cells fabricated with varying UV O 3 treatment time on NiO HTLs. Any UV O 3 treatment on NiO decreased the device performance compared to devices without any treatment. High binding energy cut off in the XPS survey scan of the freshly prepared NiO film with that of a NiO film after 10 minu tes of UV O 3 treatment on freshly prepared single crystal Si substrates cleaved from the same wafer. The data show that there is a work function shift of more than 1 eV with UV O 3 treatment. This is cannot be directly correlated to the ~0.3 V change in Voc of the solar cells. Data were corrected for sample charging by referencing the adventitious C 1s peak to 284.8 eV. ................................ ................................ .................. 77 2 12 Illum inated current density voltage ( J V ) characteristics of solar cells comprising NiO or PEDOT:PSS HTLs. The solar cells with NiO as the HTL outperform those with PEDOT:PSS due to a higher fill factor and short circuit current. .............................. 77 2 13 The plot of diode ideality factor vs. voltage for both types of solar cells. ......................... 78 2 14 Built in voltage of solar cells with either NiO or PEDOT:PSS as the HTL determined by electroabsorption. The built in voltage of the solar cells differs by the same 10 mV by which the V oc differs, suggesting the V oc is limited by the built in voltage in these solar cells. ................................ ................................ ................................ ................. 78 2 15 E xternal quantum efficiency of solar cells fabricated in this study. The EQE of solar cells with NiO is greater at nearly all wavelengths of incident light. The spectrum shift between the device is caused by optical resonance changes within the solar cell. .... 79 2 16 Simulated spectral absorption for solar cells with NiO or PEDOT:PSS HTLs. The spectra are in close agreement with the experimental spectra, indicating that the shift is due to optical resonance in the solar cells. ................................ ................................ ..... 79 2 17 Water contact angles for NiO and PEDOT:PSS HTLs. The average contact angles were 29.3 2.8 for NiO and 12.5 1.4 for PEDOT:PSS. The NiO HTL is less

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13 hydrophilic, promoting favorable wetting during solution processing and BHJ physical contact. ................................ ................................ ................................ ................. 80 2 18 AFM roughness images of 10nm thick BHJ films on NiO and PEDOT:PSS showing the contrast in physical formation of the films on different HTLs. The active layer films deposited on NiO are smoother than those deposited on PEDOT:PSS. AFM phase images of thin polymer/fullerene blend layers on NiO and PEDOT:PSS, showing a drastic change in material distribution near the HTL/active layer interface, caused by surface energy differences of NiO and PEDOT:PSS. The scale bar is 400nm and the scan area is 2 m x 2 m. ................................ ................................ ........... 80 2 19 AFM roughness images of the full ~100 nm thick BHJ films on NiO and PEDOT:PSS. AFM phase images of the full ~100 nm thick BHJ films on NiO and PEDOT:PSS. The images show nearly identical roughness and phase mapping. The scale bar is 400nm and the scan area is 2 m x 2 m. ................................ ......................... 81 2 20 Normalized device characteristics as a function of air exposure time for PEDOT:PSS and NiO based devices. Devices made with NiO HTLs are more stable when stored in ambient conditions, indicating that the NiO HTL is more stable than the P EDOT:PSS HTL in these solar cells. ................................ ................................ 82 3 1 Device structure and energy levels with respect to vacuum of the materials used in the photovoltaic cells. ................................ ................................ ................................ ........ 97 3 2 Current voltage characteristics of solar cells fabricated with Precursor A for NiO, heated to 185 o C, then UV Ozone treated to complete the formation of NiO. A drastic increase in all device performance parameters is observed after UV Ozone treatment of the heated films. ................................ ................................ ............................. 98 3 3 External quantum efficiency spectrum of solar cells fabricated with Precursor A for NiO, heated to 185 o C, then UV Ozone treated to complete the formation of NiO. ......... 98 3 4 Current voltage characteristics of solar cells fabricated with Precursor A for NiO, heated to 150 o C, then UV Ozone treated to complete the formation of NiO. At such a low temperature, the precursor was insufficiently decomposed and Ni was hardly oxidized. The resulting solar cells show negligible photocurrent and photovoltage output. ................................ ................................ ................................ ................................ 99 3 5 High resolution mass spectrum showing the isotopic distribution for the [Ni(en) 2 (OAc)] + ion. The main peak for Ni 58 is located at m/z = 237.0786. .................. 99 3 6 Absorption spectrum showing the progression of the NiO film fabrication by the low temperature route for Precursor B. ................................ ................................ ................... 100 3 7 Current voltage characteristics of the solar cells with NiO fabricated from Precursor B with the low temperature process. ................................ ................................ ................ 100

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14 3 8 External quantum efficiency spectrum of the solar cells with NiO fabricated from Precursor B with the low temperature process. ................................ ............................... 101 3 9 High resolution deconvolution of the Ni 2p 3/2 XPS spectrum before and after UV Ozone treatment. The envelope shows four peaks, one for NiO (dark red), o ne for Ni OH bonding, precursor coordination, or defect induced Ni 3+ (blue), one for NiOOH (light green), and a broad shake up peak at the high binding energy (blueish green). After UV Ozone treatment, the NiO and NiOOH peak increase, signifying the formati on of NiO and NiOOH, with a concurrent decrease in Ni OH/precursor coordination peak. ................................ ................................ ................................ ........... 102 3 10 High resolution deconvolution of the O 1s XPS spectrum before and after UV Ozone treatment. The envelope shows three peaks, one for NiO (dark red), one for Ni OH bonding/precursor coordination/defect induced Ni 3+ (blue), and one for NiOOH (blueish green at 532.3). After UV Ozone treatment the NiO and NiOOH peak increase, signifying the formation of NiO and NiOOH, with a concurrent decrease in Ni OH/precursor coordination peak. Additionally, a water peak appears after UV Ozone treatment (blue green at 534 eV). After UV Ozone treatment, the NiOOH peak is shown in bright green. ................................ ................................ ......................... 103 3 11 X ray reflectivity (XRR) spectrum of the precursor films before and after UV Ozone treatment. Curve fitting reveals that before UV Ozone treatment, there exists only one layer, while after UV Ozone treatment, there exists two layers one is the highly porous and defective NiO with a top layer of water, confirming the contribution in the XPS spectrum. ................................ ................................ ................... 104 3 12 Grazing incidence X ray diffraction spectrum showing a completely amorphous NiO film after the UV Ozone t reatment. ................................ ................................ ................. 105 3 13 Current voltage characteristics of the solar cells fabricated with the low temperature NiO process compared to those of solar cells fabricated with PEDOT:PSS or high temperature NiO. ................................ ................................ ................................ .............. 105 3 14 External quantum efficiency spectrum of the solar cells fabricated with the low temperature NiO process compared to those of solar cells fabricated with PEDOT:PSS or high temperature NiO. ................................ ................................ ........... 106 3 15 Electroabsorption measurement showing a built in potential of 1.2 eV in the solar cells with the low temperature processed NiO HTL. The large built in potential is due to the NiOOH surface dipole created during the UV Oz one treatment formation of the NiO film. ................................ ................................ ................................ ................ 106 3 16 Air stability of the solar cells with low temperature processed NiO compared to devices with PEDOT:PSS and high temperature processed NiO as references. The devices with low temperature processed NiO are slightly more air stable than those with PEDOT:PSS but less stable than those with high temperature processed NiO. ...... 107

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15 3 17 Photograph of the solar cells fabricated with low temperature NiO on PET substrates. Photo courtesy of Jesse Manders ................................ ................................ .................... 108 3 18 Current voltage characteristics of the solar cells fabricated on PET substrates. The devices achieved an average PCE of 3.7%, promi sing for further development. ............ 108 4 1 Typical absorption spectrum of PbS QDs synthesized with the hot colloidal injection technique. The same batch stored in air was measured three weeks later and showed an unchanged absorption spectrum. This indicates the QDs are well capped by the oleic acid ligands. ................................ ................................ ................................ ............. 123 4 2 Transmission electron micrographs of PbS QDs used in the photodetectors. The scale bar is 20 nm on the main image and 10 nm on the inset. The image shows nearly spherical quantum dots with approxi mately 2 nm interdot spacing, corresponding to the oleic acid ligand length. ................................ ................................ ............................. 123 4 3 Schematic of the change in spacing between q uantum dots upon exchanging oleic acid for 1,3 Benzenedithiol. The shorter benzenedithiol ligands decrease the spacing between quantum dots. ................................ ................................ ................................ ..... 124 4 4 Transmission electron micrographs of 6 nm ZnO nanoparticles used in the PbS photodetectors. The scale bar is 50 nm on the main image and 5 nm on the inset. ......... 124 4 5 X ray diffraction spectrum of the ZnO nanoparticles used in the PbS photodetectors. The diffraction spectrum confirms that the nanoparticles are 6 nm in diameter and have the Wurtzite crystal structure. ................................ ................................ ................. 125 4 6 Optical absorption spectrum of the ZnO nanoparticles used in the PbS photodetectors. The absorption edge is 365 nm, corr esponding to a bandgap of 3.4 eV. ................................ ................................ ................................ ................................ .... 125 4 7 NiO diffraction spectrum showing polycrystalline NiO in the typical rocksalt crystal str ucture. ................................ ................................ ................................ ........................... 126 4 8 Transmission spectrum of NiO films used in the PbS QD photodetector. The transmission spectrum shows that the film s are greater than 95% transparent through the entire visible and NIR spectrum. ................................ ................................ ................ 126 4 9 The physical device structure for the photodetectors. ................................ ..................... 127 4 10 Energy band diagram for the photodetector. The energy band alignments create a favorable double heterojun ction structure. ................................ ................................ ...... 127 4 11 Plot of responsivity and EQE for the photodetectors. EQE in these devices reached 24% at 1130nm and 52% at 575nm while the responsivity reached 0.2 A/W at 1130nm and 0.25 A/W at 600nm. ................................ ................................ .................... 128 4 12 Temporal response of the photodetector for light intensity of 54 mW/cm2 at 410 nm. .. 128

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16 4 13 Speed and bandwidth as a function of applied reverse bias. The bandwidth increases with applied bias. ................................ ................................ ................................ ............. 129 4 14 Photodetector bandwidth as a function of incident light intensity at 1 V applied bias. As light intensity increases, photoexcited carrier density becomes sufficient to fill traps, allowing an increase in bandwidth. ................................ ................................ ........ 129 4 15 Noise current spectral density as a function of frequency. At high applied biases, 1/f noise appeared with low f c values. As bias is decreased, J d decreased and 1/f noise is pushed to lower frequencies than measured in our experiment. At 1 V applied bias, the whole bandwidth is shot noise dominated. ................................ ................................ 130 4 16 Dark current density vs. applied bias. This plot shows extremely low dark current enabled by the novel double heterostructure. ................................ ................................ .. 130 4 17 Noise Equivalent Power and Specific Detectivity of the photodet ectors at 1 V. NEP D* was calculated to be as high as 1.1x10 12 cm Hz 1/2 /W at 1135nm and 1.2x10 12 cm Hz 1/2 /W at 600nm. For reference, reported maximum values for D* in commercial ph otodetectors are indicated. ................................ ................................ ........................... 131 4 18 Linear dynamic range (LDR) of the photodetectors measured at 1 V. The total LDR is 67 dB, co mparable to InGaAs, and sufficient for high contrast imaging. .................... 131 4 19 ambient conditions without encapsulation. Details are discussed in the text. ................. 132

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17 Abstract of Dissertation Presented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy NANOSTRUCTURED M ATERIALS FOR PHOTOVOLTAICS AND INFRARED DETECTION By Jesse Robert Manders May 2014 Chair: Franky So Major: Materials Science and Engineering Organic semiconductors and inorganic quantum dots are becoming topics of intense research due to their promising optical and electronic properties as well as their amenability to large scale high throughput solution based processing methods. Organic solar cells are rapidly improving in perfor mance with the promise of lightweight, functional and aesthetically pleasing applications. Inorganic quantum dots are especially interesting materials for next generation optoelectronic sensors, particularly in the infrared spectrum, due to their easily tu nable bandgap and controllable transport properties. First, we developed a new understanding of the interfacial properties of solution processed NiO hole transport layers (HTLs) in polymer solar cells. Polymer solar cells were fabricated with NiO HTLs with a power conversion efficiency (PCE) of 7.8%, a 14 % improvement over the reference devices. The improvement is due to an optical resonance shift in the solar cells, and an improvement in the morphology of the photoactive layer near the HTL/active layer int erface that led to increased charge extraction over the reference devices. Additionally, the solar cells with NiO were more stable in air than the control devices. In the next section, we developed a new low temperature fabrication process to incorporate s olution processed NiO HTLs into polymer solar cells. By pre heating the NiO

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18 precursor films to a low temperature, then oxidizing the films, we formed NiO at temperatures low enough to be compatible with plastic substrates envisioned for roll to roll (R2R) processing. Solar cells fabricated on standard glass substrates showed comparable performance to the community standard devices and were slightly more stable in air. As a proof of concept, solar cells were fabricated on plastic substrates using this new me thod of fabricating NiO. The cells showed favorable performance, comparable with the community standard polymer solar cells typically fabricated on glass. In the final section, we demonstrated, for the first time, an all solution processed, all inorganic m ultispectral photodetector. The devices utilized solution derived NiO and colloidal ZnO nanoparticles as charge blocking layers, and PbS quantum dots as the photoactive layer. Due to the wide bandgaps and favorable charge blocking ability of the oxides, th e photodetectors produced extremely low dark current and noise levels. They also showed large linear dynamic range, adequate bandwidth for imaging applications, and a specific detectivity comparable to, or better than, some commercially available photodete ctors. Additionally, because of the double heterojunction formed by oxides, the devices were very stable in air, remaining unchanged for five months.

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19 CHAPTER 1 INTRODUCTION 1.1 Organic Semiconductors 1.1.1 What Are Organic Semiconductors ? Inorganic semiconductors like silicon, gallium arsenide, and indium gallium arsenide dominate the modern commercial electronics landscape. An unfathomable amount of our daily life depends on these materials for our personal computers, mobile phones, busine ss machines, telecommunications, transportation, and security. However, as time marches on, innovation demands a change in the paradigm. Many processes involved in the manufacture of inorganic semiconductors are slow and expensive, and only have the capabi lity to make small area devices economically. The devices themselves, for example silicon solar panels, are bulky, heavy, and often viewed as an eyesore. As we move into the future, new applications demand thinner, lighter, flexible, more mobile, and inexp ensive devices. Organic semiconductors are promising materials for these next generation applications. Organic chemicals are general defined as any chemical containing carbon, with notable exceptions like steel and other carbides. These chemicals can be as small as the five atom methane molecule to as large as biomolecules like DNA which can contain billions of atoms. The structure of the three main types of organics small molecules, polymers, and biomolecules, are shown in Figure 1 1 What all of these molecules have in common is that the constituent atoms are bound with covalent bonds and the molecules are held together solid form predominantly by weak van der Waals forces Because of these molecular separation in the solid phase and weak van der Waals forces, the bulk properties often resemble those of the individual molecules Organics are often flexible and lightweight, making them attractive for many applications where ergonomics and economics are important. Because the number of organic

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20 molecules we can synthesize is essentially infinite, the optical and electronic properties of these materials can be tuned with atomic precision. Organic semiconductors, in particular, are materials which contain conjugated systems. A conjugated system is a series of overlapping p orbitals across an intervening sigma bond that allows for delocalization of p orbital electrons, analogous to that of delocalization in inorganic semiconductors. This will be discussed more in section 1.1.2. The two main classes of organic se miconductors are small molecules and polymers. Small molecules contain tens or hundreds of atoms, while polymers are long chains of bonded monomer units that can have a molecular weight of millions of Daltons. With the discovery of conducting polymers [1] organic electronics has steadily become a large active area of academic research and commercialization. Earl y devices based on organic semiconductors were unstable and inefficient [2 5] However, much progress has been made through s ynthetic chemistry and device physics advancement. In particular, organic light emitting diodes (OLEDs) have gone from newly practical in 1987 [4] to commercially available in a large number of mobile phones and upcoming televisions in 2014. In recognition of the magnitude of the advancement in organic electronics, Alan Heeger, Alan MacDiarmid, and Hideki Shirakawa were awarded Some major advantages organics have over traditional inorganics is that they can be deposited in devices from solution, or low temperature sublimation. The ma terials do not need to adhere to strict lattice matching constraints and be defect free. Furthermore, because of the tight localization of charges in organic semiconductors compared to inorganic crystals, even amorphous materials can make highly efficient devices like solar cells [6] since defect structures

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21 do not affect transport and absorption as strongly. In order to tailor organic molecules for next generation devices, detailed knowledge o f their electronic structure must be understood. 1.1.2 Electronic Structure of Organic Semiconductors Organic semiconductors are characterized by their electronic structure being fundamentally different from that of inorganic crystals. In inorganic semicon ductors, when a sufficient amount of atoms bond to form a bulk structure, the atomic orbitals must shift to different energies as governed by the Pauli Exclusion Principle [7] In shifting the large number of atomic orbitals effectively merge into continuous bands of energy levels. In this bulk structure, the wavefunctions of electrons and holes are highly delocalized. However, in organic semicond uctors, the action of forming energy levels is quite different. 1.1.2.1 Molecular orbital c onstruction Organ ic semiconductors owe their conductive and optical properties to being conjugated systems. As described above, a conjugated system is a series of overlapping p orbitals across an intervening sigma bond that allows for delocalization of p orbital electrons, analogous to that of delocalization in inorganic semiconductors. Thus conjugated systems a re composed of alternating single and double bonds where the double bonds are actually smeared across the entire conjugation length Single bonds are composed of bonds, while double bonds are composed of both and bonds. Along the conjugation length carbon is sp 2 hybridized, meaning the atomic s orbital and two of the three atomic p orbitals have merged to form three hybrid sp 2 orbitals. T he p z orbital remains unhybridized and extends above and below the ring plane that is formed by the bond of over lapping sp 2 orbitals. Above and below the C C bond, p z orbitals overlap edge on to form bonds. A depiction of the sp 2 and p z orbitals and their energy level construction in this situation for two carbon atoms (without conjugation) is shown

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22 in Figure 1 2 In larger systems, e.g. benzene pentacene, polyphenylene vinylene, etc. conjugation can occ ur since there are more than four carbon atoms with alternating single and double bonds As the number of carbon atoms (and thus conjugation length) increases the bonding and antibonding orbitals must become degenerate following the Pauli Exclusion Principle in the same manner as inorganics. This leads to small band formation, with the band edge of the orbital band being the highest occupied molecular o rbital (HOMO) and the band edge for the orbital band being the lowest unoccupied molecular orbit (LUMO ) The HOMO and LUMO are thus analogous to the valence band edge and conduction band edge of inorganic semiconductors, respectively. In physical space, this means that the bonds are smeared across the entire conjugation length and the electrons become delocalized along the conjugation length. Even though there is delocalization of electrons in p conjugated systems, the spatial extent of the delocalizat ion is typically far smaller than that of inorganic semiconductors, particular for small organic molecules and low molecular weight polymers. This delocalization is epitomized by aromaticity of a n isolated benzene ring, as shown in Figure 1 3 In organic s emiconductor devices, the HOMO LUMO gap, or transition is critical and usually lies between 1 4 eV. With the construction of the energy levels of organic semiconductors now known, we must understand the nature of the excitation and motion of the electrons and holes within this structure. 1.1. 2.2 Excitons An exciton is a Coulomb correlated (bound) electron hole pair. Thus, an exciton is a chargeless quasiparticle that exists in solids in the excited state. It can also be viewed as an excited state of a molecule [8] Thus, an exciton is a way to transport charges without generating current. The formation of excitons occurs either af ter injected holes and electrons meet, or after

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23 the material is photoexcited by a photon with energy larger than the bandgap (or HOMO LUMO gap). In the case of the photovoltaics and photodetectors discussed in this dissertation, excitons are generated by p hotoexcitation of the photoactive materials. There are three types of excitons the Mott Wannier exciton, Frenkel exciton, and charge transfer (CT) exciton. These three ex citons are depicted in Figure 1 4 The Mott Wannier exciton is characterized by a la rge radius, many times greater than the lattice constant of the material. This exciton is mostly found in inorganics where the dielectric constant is high so the screening between the electron and hole is small, and thus their delocalization is large. Thes e excitons typically have very low binding energies, on the order of 10 meV, so they are dissociated into an individual electron and hole quite easily. The Frenkel exciton is typical of organic semiconductors and is characterized by a large binding energy up to ~1 eV. The Frenkel exciton has a small radius due to the low dielectric constant of the semiconducting medium. The spatial extent of the exciton is limited one unit cell (or molecule in organics). Because of the low dielectric constant of organics an d subsequent large binding energy of Frenkel excitons, dissociation into free charge carriers requires clever engineering techniques. In modern donor acceptor polymers, where an electron donating unit alternates with an electron accepting unit on the backb one [9] Frenkel excitons form between the donor and acceptor monomers along the backbone of the polymer. Thus a dipole moment is induced along the polymer backbone. This dipole moment is correlated with the exciton binding ene rgy and ultimately the ability of the exciton to be dissociated into free carriers [8] In this vein, chemical structure plays a large r ole in charge separation efficiency, and ultimately on device efficiency.

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24 The charge transfer (CT) exciton is an exciton that extends over two or three lattice sites (molecules). One carrier is localized on one molecule while the other carrier in the exciton is localized on an adjacent molecule. CT state energy levels are quite complex and binding energies are highly dependent on the materials system [10,11] This exciton is particularly important in determining the performance in solar cells and can signify pathways to more efficient devices. 1.1.2.3 Charge t ransport In inorganic semiconductors, charge carriers are highly delocalized (mean free paths are much larger than the lattice constant) and the transport bandwidths are relatively large. Thus, band transport is the dominant mechanism in crystalline inorganic semiconductors and carrier mobilities are often greater than 1 cm 2 /V s [7] The picture is quite different in organic semiconductors. In organics, charge transport is dominated by polaron hopping. A polaron is a qua siparticle made of a charge carrier and its local lattice/molecule distortion caused by the serves to reduce its mobility and increase its effective mass. I n this model, the distortion of the local environment creates a potential well out which of the carrier must hop to continue m oving through the material. With each hop, the carrier distorts its new local environment, essentially creating its own potential well always hopping between new potential wells This is depicted in Figure 1 5 Therefore, the charge carriers are highly localized (mean free paths the size of one lattice constant /monomer or smaller) on individual molecules or polymer segments that are held together by weak van der Waals forces. As a result of this localization, carriers cannot easily flow through the material and must hop between molecules. While single crystals of organic molecules can have mobilities greater than 1 cm 2 /V s [12] the more common amorphous organics suffer from much lower mobilities, often lower than 10 3 cm 2 /V s

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25 Because of the large number of organic molecules, the ir different molecular arrangements and subsequent charge transport mechanisms, energetic disorder, and charge localization details, several charge transport models exist. However, the most commonly applied to fit current voltage (J V) data come from space charge limited conduction (SCLC). In particular, the Mott Gurney law, while only valid under certain conditions (1. Single carrier conduction, 2. Trap free material, 3. Negligible intrinsic carrier concentration, and 4 Parallel plate electrodes), is ofte n employed as a starting point for analysis. The Mott Gurney law can derived from solving the Poisson equation and the continuity equation [13] with appropriate boundary conditions or derived from geometrical considerations [14] and is written as (1 1) w here J is the current density, is the mobility, is the dielectric permittivity F is the electric field, and d is the thickness between the electrodes. This equation, and va riations such as that which includes field dependent mobility [15] can be us eful in determining SCLC carrier mobility in new materials as long as injection contact is ohmic. In this section we have constructed organic semiconductors from an atomistic perspective, developed photoexcitation dynamics, and charge transport phenomena. This knowledge is crucial to developing useful devices, such as the photovoltaic cells presented in this dissertation. 1.2 Organic Photovoltaics 1.2.1 Background on Organic Photovoltaics In a society concerned with environmental, economic, and geopolitica l consequences of energy consumption, new energy sources are needed to secure a prosperous future. Organic photovoltaics (OPVs) are emerging as viable alternative energy sources for several important

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26 reasons: they are compatible with low cost, large scale manufacturing, can easily be integrated into urban or rural environments, and can be manufactured on flexible substrates with a wide variety of shapes and sizes, making artistic, yet functional applications realistic. Since the first organic photovoltaic d iscovery [3] and first reali stic organic solar cell [5] much has been learned about the photophysics and chemistry in these devices. Recent advances in the understanding of the chemistry and physics of OPVs are pushing the energy harvesting efficiency, or power conversion efficiency, to new heights. Record high power conversion efficiencies in solar cells compatible with large scale roll to roll (i.e. newspaper like) printing reaching beyond 8% have been demonstrated [6,16] To bring OPVs to the forefront of the renewable energy push, continuous development of new materials and processing techniques in around the world is essential. As the world increasingly focuses on providing sustainability to future generations, these recent advances and ongoing development projects will certainly play an integral role in solving our global energy problems. 1.2.2 Measurement and Fundamental Electrical Characteristics same standards, lamps which generate the Air Mass 1.5G (AM 1.5G) solar spectrum are used The spectrum is shown in Figure 1 6 This spectrum is defined as the solar spectrum of incident light on the surface of the E arth at a solar zenith angle of 48.2 o c orresponding to an effective atmosphere thickness of 1.5 times actual atmosphere thickness The fundamental measurement of solar cells is the current density voltage ( J V ) measurement. For this, the AM 1.5G spectrum at 1000 mW/cm 2 irradiance is incident on the solar cell while the voltage is swept and the output current is measured. A typical J V plot is shown in Figure 1 7 Several important parameters are extracted from the J V plot:

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27 1) Open circuit voltage ( V oc ) the voltage output of the solar cell whe n the solar cell is placed in an open circuit. This is the maximum output voltage of the solar cell run in photovoltaic, or power generating, mode (quadrant III). 2) Short circuit current density ( J sc ) the maximum current output of the solar cell in powe r generating mode, equivalent to the current generated when the device is placed in a shorted circuit connected only with itself. This is usually reported in mA/cm 2 4 ) Maximum power point ( P max ) The maximum power output, found experimentally by multiplying each measured J by its corresponding V value in quadrant III of the J V plot. When extracting P max from the J V plot, P max is the maximum power density in W/cm 2 or mW/cm 2 5) Fill factor (FF) This is a parameter which demonstrates the ratio of actual maximum output to theoretical maximum output for the solar cell. It is found numerically by (1 2 ) 6) Power conversion efficiency (PCE, ) This is the percent of opt ical power incident on the solar cell that is converted into electrical power. It is calculated by (1 3) where P incident is the incident power in the AM 1.5G spectrum (1000 mW/cm 2 ). This is the efficiency number that ultimately defines how desirable a solar cell is. 7) External quantum efficiency (EQE) This quantity is defined as the ratio of electrons extracted from the solar cell per second to the number of photons incident on th e solar cell per second. This is measu red as a function of wavelength, and thus is a good measure of how

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28 well the solar cell is performing at specific wavelengths. It is represented algebraically by the relation (1 4 ) w light, J photo ( ) is the photocurrent generated by the device as a function of wavelength, and P incident ( ) is the incident power at the measured wavelength. Typically, the J sc should be cross checked with the EQE by integrating the EQE with the AM 1.5G spectrum to calculate J sc as (1 5) where E in ( ) is the spectral irradiance in the AM 1.5G spectrum. Now that the critical parameters of solar cell performance are understood, we must develop the device structures to show how real solar cells are constructed. Then, the limitations on achieving these pa rameters and opportunities for advancement will be understood. 1.2.3 Organic Solar Cell Device Structures The first practical organic solar cells were made in a bilayer planar heterojunction structure [5] In this structure, an electron donating and electron accepting material are sandwiche d between the cathode and anode without any donor acceptor interfacial mixing. In this structure, for example, light is absorbed in the electron donating species which gener ates an exciton. The exciton then diffuses to the planar heterojunction interface between the electron donor (donor) and electron acceptor (acceptor), where the energy level offset between the two materials provides the energy needed to dissociate the exci ton into free carriers. In organic materials, as a result of the strong localization discussed previously, the exciton diffusion length is typically only on the order of a few nanomet ers This short diffusion length poses a serious problem for

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29 the use of a bilayer architecture due to an absorption dissociation trade off. In order to generate a large amount of current, the photoactive layer has to be thick enough to absorb a large amount of photons. However, in making the photoactive layer very thick, the di stance an exciton must diffuse to reach the donor acceptor interface to dissociate into usable free carriers is increased well beyond the exciton diffusion length. Th is creates parasitic recombination which hinders the performance of the cell. To solve thi s issue, the bulk heterojunction (BHJ) was created [17,18] A bulk heterojunction is the mixing of the donor a nd acceptor material into a single layer, where they form small interconnected domains throughout the thickness of the active layer. These small domains should ideally be no more than five to ten nanometers thick so as not to exceed the exciton diffusion l ength in the material. The domain size and distribution and film roughness (together termed the morphology) can be controlled by using different solvents and annealing conditions [19,20] These two device structures are contrasted in Figure 1 8 In polymer solar cells, the only type c onsidered in this dissertation, the donor material is a semiconducting polymer, while the acceptor material is typically some functionalized fullerene derivative, such as PC 60 BM or PC 70 BM ( Appendix A ). The functionalization with a side chain gives the full erene its desirable solubility so it can be codeposited with the polymer from one stock solution by spincasting or spray coating One issue in this structure is the vertical phase orientation between the donor and acceptor. Since the donor and acceptor are intimately mixed, each material is in physical contact with each electrode. This drastically reduces the shunt resistance, potentially lowering the fill factor and open circuit voltage of the cell. To solve this issue, interfacial layers are inserted betw een the contacts and the BHJ photoactive layer to block the wrong carriers from reaching the electrode Selecting and tailoring materials to obtain the necessary interfacial energetics and surface properties is critical to fabricating an efficient device.

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30 1.2.4 The Importance of Interfaces In order to prevent parasitic current leakage, BHJ solar cells employ interfacial layers. In the conventional solar cell architecture, a hole extraction/hole transport (HEL/HTL) layer is inserted between the anode and the BHJ layer to block electrons from being extracted at the anode. Traditionally, this HTL has been poly(3,4 ethylenedioxythiphene) doped with poly(styrenesulfonate) (PEDOT:PSS), a p type conductive polymer. Its work function of 5.2 eV is usually well suited for conventional solar cells. However, its inability to effectively block electrons, its acidity, and its sensitivity to moisture all counter its favorable transport characteristics [21] Additionally, many new donor polymers have very low lying HOMO levels [22] [23] ; this necessitates e ven deeper Fermi levels in hole transport layers for favorable energy level alignment to prevent barriers to carrier extraction. Materials like NiO and MoO 3 have been used successfully as large work function interfacial layers [6,16,24 26] Along with Fermi level alignment, the surface energy of these transport materials are critical parameters that needs to be tailored to create an effective mating to the BHJ layer. The surface energy of the hole transport layer is critical in conventional polymer solar cells. If the surface energy of the HTL is not appropriate for the BHJ layer being deposited on top, the interface between the H TL and the BHJ will be rough, which will lead to interfacial recombination and a decrease in shunt resistance of the solar cell. Different methods to tailor the surfac e energy of these materials can be used, including UV ozone treatment, oxygen plasma, and depositing self assembled monolayers to create a surface dipole. These surface treatments can not only alter the surface energy, but alter the effective work function since the surface dipole created by these t reatments will increase or decrease the barrier across which the electrons or holes travel. The ideal situation would be to select a material for the HTL that does not need any surface modification to function optimally. Chapter 2 of this dissert ation describes such a

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31 case, while C hapter 3 of this dissertation describes a case where using a surface treatment can actually be beneficial to creating polymer solar cells on the path to being compatible with large scale, high throughput processing. 1.3 Inorganic Nanocrystals and Quantum Dots 1.3.1 Fundamentals of Nanocrystals and Quantum Dots Nanocrystals and quantum dots are tiny groups of atoms ranging in size from a few nanometers to tens of nanometers, containing several hundred to several thousands of atoms. This small size endows them with optical and el ectronic properties between those of molecules and the bulk. Because of these interesting properties, there are many applications in development and certainly many unseen at this time that will utili ze these unique materials. Nanocrystals can be pure metals, metal alloys, semiconductors, and magnetic materials and can be used in biological imaging [27] optical sensors [28] light emitting diodes [29] and photovoltaic cells [30,31] Strictly speaking, quantum dots ar e a subset of nanocrystals whose electronic and optical properties are modulated through the quantum confinement effect [32 34] In this dissertation, all of the real in chapter 5 are in fact semiconducting quantum dots. The distinction in this work is merely for convenience for the work presented in chapter 5 Quantum dots are particularly interesting because changing the size and shape of the material can drastically alter many of their properties. Among the most important tunable properties for optoelectronic devices are t he optical absorption and emission profiles, optical bandgap, an d charge transport properties For example, PbS quantum dots can be tuned to absorb light across the entire near infrared (NIR) spectrum, with an absorption tail that covers the entire visible spectrum [35] and CdSe quantum dots can be turned to emit light across the entire visible spectrum, leaving the future open for quantum dot displays and solid state lighting [29]

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32 Another reason working with quantum dots and other nanocrystals is attractive to researchers is that the y are relatively easy to synthes ize and the produ cts are often bulk defect free. One disadvantage of quantum dots is their propensity to agglomerate and lose their quantum properties without proper surface modification. This agglomeration is caused by the h igh surface to volume ratio of these materials, making them thermodynamically unstable. Through this disadvantage comes opportunity, however. Researchers are free to modify the surfaces by synthesizing the quantum dots with various ligands which can later be exchanged to crosslink multiple quantum dots, change the polarity of solvent used in processing, and improve the colloidal stability [36 39] Thus, th ese materials are readily compatible with multiple processing conditions and synthetic routes. 1.3.2 Synthesis of Quantum Dots and Nanocrystals There are many different routes t o synthesizing nanocrystals Among them are gas phase, solid phase, and liqui d phase. Liquid phase synthesis has advantages over gas phase and epitaxial growth mechanisms. Gas phase growth can be performed by flame, plasma, laser mediated and aerosol growth These routes are attractive due to the lack of vacuum constraints and ine xpensive processing in many of the gas routes, but typically results in high polydispersity and may not be easily compatible with many existing optoelectronic device fabrication schemes [40] Solid phase growth is typically performed in metal organic chemical vapor deposition (MOCVD) or molecular beam epitaxy (MBE ) reactors [41] MOCVD and MBE reactors are often a bottleneck to high throughput and inexpensive synthesis and fabrication due to the vacuum and lattice matching requirements of the materials involved. Because of these drawbacks, the most common, and the method used in this dissertation, is the liquid phase synthesis. Within the liquid phase synthesis route, many techniques exist, including the reverse micelle [42] microemulsions [43] and hot colloidal injection [38] The technique which has garnered

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33 the most progress for optoelec tronic materials is the hot colloidal injection technique because of the flexibility of reaction parameters and wide range of achievable products. There are many reaction parameters that can have an effect on the shape and size of the quantum dots that are produced. Among them are reaction temperature, time, precursor concentration ratio, and reaction vessel pressure. Typical growth conditions prescribe reaction times on the order of a few minutes and temperatures between 100 400 o C for different materials; lead chalcogenides are typically grown at lower temperatures than cadmium chalcogenides, for example. Quantum dots synthesized with the hot colloidal injection technique have been used in devices such as photodetectors [28,44] photovoltaic cells [31] and light emitting diodes [29,45] spanning the ultraviolet (UV), visible (VIS), near infrare d (NIR), short wave infrared (SWIR), and the mid wave infrared (MWIR) spectra. In this dissertation, growth of colloidal quantum dots is performed by the hot colloidal injection technique [38] In a general synthesis using this method, organometallic precursors are heated and dissolved in non coordinating and coordinating solvents. Upon heating, stirring and dissolution of one precursor (usually the cation), the anion precursor solut ion is injected into the cation containing solution. Quantum dot formation follows classical nucleation and growth dynamics. At the appropriate time in the reaction, the QDs are extracted and wash ed with various solvents, and dispersed in any number of sol vents for further use. The experimental setup for quantum dot synthesis is shown in Figure 1 9 The end products are nanocrystals with dimensions on the order of nanometers, surrounded by stabilizing, electrically insulating capping ligands, either organic or inorganic. As discussed above, t hese capping ligands play multiple roles. Among them are nan ocrystals shape and size control, electronic transport property modulation, electronic surface state passivation, and colloidal stability [39,46 48] These ligands can

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34 be exchanged in a post synthetic process to furth er modulate the properties of the quantum dots. To fully understand how to synthesize usage appropriate quantum dots, it is critical to understand the formation and growth mechanism of these materials. 1.3.3 Growth Mechanism of Quantum Dots and Nanocrystal s The steps involved in the entire synthesis is shown in Figure 1 10 When precursors are injected into the reaction ves sel, the precursor solution becomes supersaturated and eventually reaches a concentration known as the nucleation threshold. When this i s reached nucleation of the precursors occurs to partially relieve supersaturation by lower ing the concentration in solution. During nuclea tion, small amounts of atoms are bound together in clusters and the growth of these clusters is governed by free ene rgy constraints Nucleation and growth of quantum dots follow classical nucleation and growth theory where minimizing the Gibbs free energy ( G) of the system is the ultimate goal [38,49] The total Gibbs free energy of the nucleating system ( G) can be thought of as the sum of the volume free energy ( G v ) the free energy difference between the solid phase and liquid phase, and the surface free energy ( ) the energy associated with the formation of the interface between the liquid and solid phase s during nucleation. Thus, the total free energy is given by, (1 6 ) where the familiar formulas for vol ume and surface area of a spherical particle appear with the associated free energies. The volume free energy is negative, while the surface free energy is and solid makes it thermodynamically unfavorable, while creating the bulk of the material to relieve supersaturatio n is favorable. Analyzing these conditions further, we find that there is

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35 some critical radius r*, where if the cluster of atoms stops growing at a size smaller than r*, it will re dissolve and be lost as the way to minimize the free energy This is because the energy cost of creating the surface is not overcome by the reduction in free energy of creating the new small solid volume If the cluster grows to be the size of r*, then it is able to minimize free energy by continuing to grow. Thus, r* can be thoug ht of as the minimum nucleus size. This scenario is shown graphically in Figure 1 11 The critical radius is then easily fou nd by setting the deriv ative Eq uation 1 6 to zero and solving for r. The result yields the following expression for r*: (1 7) A few assumption go into applying classical nucleation and growth dynamics to quantum dot synthesis: 1) we assume that all QD clusters and nuclei are spherical, 2) we neglect the energetic difference between different crystal faces and crystal str uctures, and 3) we neglect ligand/surfactant interaction with the nucleus during the growth process. Once enough nucleation events have occurred to lower the precursor concentration below the nucleation threshold, nucleation stops and the reaction proceeds through the growth process [38,50] However, the reaction solution is still supersaturated, so the growth process is predominantly limited by the reaction rate at the surface of the nuclei Once the reaction solution is no longer supersaturated, t he growth pro cess becomes limited by diffusion of the precursor solute to the surface of the nuclei instead of by the surface reaction that attaches the solute to the growing nuclei. There are some important considerations on mon odispersity that one must consider to obtain high quality materials from this synthesis. Since the growth kinetics of one nucleus is governed by the same laws as all other nuclei, it is necessary to ensure that as many nuclei as

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36 possible form at the same time to ensure a highly monodisperse sample. Thus, injection of the precursors into the reaction must be swift to ensure a rapid burst of nucleation events. Additionally, the growth rate, monodispersity, and r* are dependent on the precursor concentrations [47] Fig ure 1 12 shows the relationship between these properties. As the reaction progresses and precursor concentration is depleted, the critical size increases During this process, if the size of all the particles is larger than the critical size, then the rema ining solute will tend to grow the smaller particles more than the large ones because of the need to minimize the surface to volume ratio as described in Figure 1 11 This creates a size focusing effect and improves the monodispersity. If the opposite is t rue and the size of all the particles is lower than the critical size, the smaller particles will be sacrificed to the benefit of the larger particles by the process of Ostwald ripening. Ostwald ripening creates a highly polydisperse sample and is generall y not desirable. As stated above, it is better that the nucleation events are separated in time from the growth stage because of the complex nature of the particle growth. However, in real life reactions, there is often overlap between the nucleation perio d and growth period, generating a finite size dispersion. It is clear that precise control of the synthesis is required to obtain materials with the desired optical and electronic properties. 1.3.4 Optical and Electronic Properties Pione ering work by Efr os and Efros [32] and Brus [33,34] laid the theoretical grou ndwork of the field. In his work [34] Brus developed a relationship between the diameter of a QD and its ameter of the QD is smaller than the excitonic Bohr radius, (1 8)

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37 where a B is the excitonic Bohr radius, e is the elementary charge, and m e and m h are the effective masses of the electron and hole, respectively. In the strong confinement regime, the electronic states are nearly analogous to the 3D potential well from fundamental quantum mechanics. Energy bands are split into discrete states and the lowest conduction band state and highest valence band state are s eparated by the bulk bandgap plus the confinement energy seen in the potential well equations. In Figure 1 13 the evolution of the electronic states from a molecule to th e bulk, with QDs in the middle. From Figure 1 13 it is clear that as the number of a toms increases, the size of the quantum dot increases, and thus the energy levels tend toward the bulk, indicated by a reduction in bandgap. It is now apparent that one of the main advantages of employing QDs in any optoelectronic device is the tunability of its bandgap simply by changing the size, an easy feat, which involves control of the growth time and temperature during the reaction. Utilizing the Brus model or effective mass approximation, the bandg ap E g can be resolved from the diameter of the qu antum dots as: (1 9) where R is the diameter of the quantum dot, m e and m h are the effective masses of the electron and hole, respectively. The middle term represents the infinite potential well with its familiar 1/R 2 form, and the third term represents a screened Coulombic attraction term between the electron and hole in the nanocrystal. The Brus model has been shown to be only a rough approximation to the experimentally determined bandgap energies by many groups. To determine a better fit to experimental data, the hyperbolic band model was developed to provide a more complex band structure than the

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38 relatively simple parabolic potential well structure The relationship between the quantum d ot bandgap and its diameter for the hyperbolic band model is shown in Eq uation 1 10 below [51] : (1 10) There are still deficiencies with this model in its application to every quantum dot system and with PbS at small QD diameters in particular as real materials deviate from the computational assumptions made for all of these models. In recent work [52] experimental measurements of the bandgap and quantum dot diameter have been used to generate a modified Brus model that have provided incredible accuracy for PbS quantum dots The size of the PbS quantum dots in this dissertation is determined by transmission electron micr oscopy and confirmed by Eq uation 1 11 b elow. (1 11) The bandgap is the most important property that needs to be considered for applications like photodetectors Since the bandgap of these materials can be tu ned by changing their size, the absorption spectrum and t hus the working spectr um of devices can also be tuned Figure 1 14 shows the tunability of PbSe quantum dots used in quantum dot based photodetectors across the entire NIR spectrum just by changing their size. This tunability, combined with large scale sol ution processing opens up many avenues for production of interesting optoelectronic devices. To develop high performance optoelectronic devices with these materials, it is essential to understand their charge transport characteristics. 1.3.5 Charge Transpo rt and Excitons in Quantum Dots Studies are still being performed to elucidate the exact charge transport properties in quantum dots and quantum dot films. However, much has already been uncovered about the

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39 nature of electrons, holes, and excitons in these unique materials. Because the size of quantum dots is between that of atoms and bulk inorganic semiconductors, many interesting transport phenomena arise. Typically band transport is seen in crystalline materials with a high degree of carrier wavefuncti on delocalization in the material, and carrier mobilities are much higher than for amorphous or disordered materials. Although semiconducting quantum dots are generally small single crystals, the transport is fundamentally different in quantum dots due to their small and discreet nature and discontinuous packing in films due to the capping ligands Since quantum dots are a few nanometers in size and physically separated from each other in solid films by capping ligands, charge transport is typically by the hopping mechanism [39] In the hopping mechanism, charges are momentarily localized in small discreet potential well by either an applied bias or thermal energy. This is shown schematical ly in Figure 1 15 In a quantum framework, this hopping is more accurately described as the delocalization and movement of the carrier wavefunction over multiple QDs. In this framework, it is clear that hopping is governed by quantum statistics and is anal ogous to quantum tunneling ; therefore, we can apply quantum tunneling from introductory quantum mechanics to obtain a relationship [39] between the tunneling or hopping rate, the energy needed to hop to the next QD, and the sp acing between QDs, as shown in Eq uation 1 12 : (1 12) where is the tunneling rate between two orbitals of QD neighbors, m* is the effective mass of the charge carrier, E is the height of the tunneling barrier, and x is the shortest distance

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40 between two quantum dots. While this is an ideal hopping mechanism without a temperature dependence, experimental conductivity data have been fit to other Arrhenius type relations [53] In quantum dot films each QD represents an individual potential well out of wh ich an electron or hole must hop to continue towards an electrode. This hopping process can be made ligands attached to the surface. Long ligands like oleic ac id a re approximately 2 nm in length and often pose a significant barrier to transport effective transport while shorter thiol ligands and atomic passivation techniques drastically reduce the inter QD spacing and improve transport and device performance [54] Photoexcitation in quantum dots shares some qualities with that in organic semiconductors [55] In both QDs and organics, photoexcitation takes place in molecular orbitals, so well known constructions of the electronic states are used to describe the process. According to fundamental molecular orbital theory, organic molecular orbitals are composed of linear combinations of overlapping atomic orbitals of the same type i.e s, p, d, f. Additionally, the conjugated bonds responsible for conduction are all made with the sam e atoms carbon (with the occasional heteroatom inserted into the conjugation) This, combined with the relatively low dielectric constant or organics, leads to highly correlated electrons and holes in highly localized excitons, as discussed previously. H owever, the case is different with quantum dots. In quantum dots, unlike atoms are bonded together, e.g. Pb and S. Thus, the molecular orbitals of these materials are composed of one atomic orbital from a cation and one from an anion. The unoccupied molecu lar orbital is a linear combination of from the cation atomic orbitals (Pb 2+ ), while the occup ied molecular orbital is a linear combination of the anion atomic orbitals (S 2 ). Therefore, photoexcitation and exciton formation is actually a charge transfer p rocess in

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41 quantum dots, where an electron is excited from the HOMO of S 2 to the LUMO of Pb 2+ Furthermore, the inorganic semiconductors that make up quantum dots have much higher dielectric constants than organics; thus, excitons are much more loosely bou nd. The spatial constraint which keeps the electron and hole close together in the exciton is the size of the quantum dot. The excitonic properties of quantum dots, coupled with the ability to tune the transport properties by relatively simple synthetic pr ocedures makes them attractive for use in next generation optoelectronic devices. In particular, the tunability of the bandgap, the tunability of the carrier mobility, and the ease of processing and synthesis makes them ideal candidates for next generation photodetectors. 1.4 Quantum Dot Photodetectors 1.4.1 Fundamentals of Photodetectors Photodetectors, like photovoltaic cells, are devices which produce and electrical signal in response to an optical input signal. Unlike photovoltaic cells, however, photod etectors are run with a bias applied to facilitate charge extraction and improve speed of response. These devices can sense light from the ultraviolet through the infrared wavelengths Photodetectors are used in many modern technologies including CCD and CMOS sensors in digital cameras, receivers in fiber optic communication systems free space communication receivers, night vision goggles, and television remote control receivers. Of particular interest in commercial technology and researchers alike are infrared photodetectors because the main fiber optic telecommunications windows in the NIR spectrum at 1.3 m and 1.55 m. Also, atmospheric windows from 3 5 m and 8 12 m are important for Traditional silicon photodetectors cover wavelengths of 300 nm to 1.1 m, and are used in commercial applications

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42 like CCD and CMOS sensors. InGaAs and Ge are commonly used for NIR sensing applications like optical communications, non thermal nig ht vision, as they cover ~800 nm to 1.7 m InGaAs, HgCdTe, bulk PbS, InAs, and InSb are commonly used for SWIR applications between 1.8 and 3 m. HgCdTe (abbreviated MCT) is the most common material used for 3 8 m MWIR and 8 12 m LWIR applications. A ch art of operating wavelengths for different inorganic photodetectors is found in Figure 1 16 Although these photodetectors cover a wide range of wavelengths in the infrared, they are often made with expensive, small area, and slow epitaxial growth techniq ues such as MBE or MOCVD that require a high degree of lattice mat ching between the materials After this fabrication, the individual sensors are assembled into an array called a focal plane array (FPA) and bump bonded onto read out integrated circuits (RO ICs) that transmit the signal to the user. A schematic of the final device fabricated using these methods is shown in Figure 1 1 7 This process is costly and only allows for small area devices. Solution processed semiconductors offer an attractive alternat ive to this process, as they can be made quickly in large quantities and deposited by fast and cheap techniques like spin coating or spray coating. Organic materials have been used in photodetectors [56] but the electronic structure of organic materials does not easily allow for optical absorption beyond ~1.3 m. Thus quantum dots, which are capable of reaching MWIR wavelengths can be easily integrated into these devices. Instead of meeting lattice matching requirements and needing bonding to the ROICs, quantum dot photodetectors could be deposited directly onto the pre pat terned ROICs using these cost effective and fast processing methods With the wide array of different materials and designs, it may not be obvious as to which type of photodetector is best for a pa rticular application. In order to compare different

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43 photodetectors and choose which one is best suited for a particular application, one must understand how these devices are characterized. 1.4.2 Performance Characteristics In order to understand the quality of a photodetector, it is necessary to understand its operating conditions and performance characteristics There are several different quantities that describe how well a photodetector operates, including those related to light absorption, noise l evels, dynamic range, and stability. The most fundamental quantity is known as the external quantum efficiency, or EQE. This is defined th e same as in photovoltaic cells as the ratio of charge carriers extracted per second to the number of photons incident on the device per second. The responsivity is a key signifier of the strength of the photoresponse of the photodetector. It is the number of amperes of current extracted from the device per watt of incident light power. In mathematical form it is can be e xpressed as (1 13) w here R is the responsivity, J photo is the photocurrent of the device, P in is the input optical power, EQE is the external quantum efficiency, e is the elementary charge, h is P is the wavelength of the incident photon, and f is the frequency of the incident photon. Th e bandwidth of a photodetector is a good measure of its speed of response to an input signal. It i s the frequency at which the output signal powe r drops to 3 dB, or half of its maximum value. It can be calculated from the photocurrent on/off transient by measuring the characteristic time constant and calculating (1 14)

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44 wher e B is the bandwidth, and is the characteristic time constant defined as follows. The time constant is time it takes for the photovoltage or photocurrent to rise to (1 1/e) times the maximum value, or ~63.2% of its maximum value when light is shone upon it. Equivalently, i f the response transient has a symmetric rise and fall, then this is equivalent to the time it takes for ~ 36.8% of the maximum value after the incident light is removed The linear dynamic range (LDR) describes the magnitude of the photoresponse to different intensities of light and is typically reported in decibels Specifically it is calculated from Eq uation 1 15 below, (1 15) where P max is the highest inc ident light power and P min is the lowest incident light power of the range in which the photodetector response is linear with the incident power. Another way of saying this is that P min and P max are the upper and lower bounds of the range of incident optic al power in which the device produces the same responsivity. In imaging photodetectors, this parameter is important because it represents the range of brightness of objects that the photodetector can image on a linear brightness scale without distorting th e true brightness. The noise equivalent power (NEP) in watts is the input signal power at which the output signal to noise ratio is unity. It represents the noise power floor of the photodetector and can be written as (1 16) where < i 2 > is the integrated mean square noise spectral density in the device bandwidth and R is the spectral responsivity.

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45 The specific detectivity ( D* takes into account the signal to noise ratio, th e spectral responsivity and the bandwidth of the detector all in one parameter. Following the definition for NEP, D* is defined as (1 17) where A is the device area, B is the bandwidth, and NEP is the noise equivalent power. Typical values of commercially available Si and InGaAs photodetectors are on the order of 10 12 to 10 13 cm Hz 1/2 /W, and Ge photo detectors are often on the order of 10 11 cmHz 1/2 /W. These performance characteristics vary drastically depending on the type of photodete ctor and its device structure. In the next section, various device structures and operating mechanisms will be contrasted. 1.4.3 Quantum Dot Photodetector Device Structures Two different device structures have been developed for quantum dot photoconductors and photodiodes. These two are structures are depicted in Figure 1 18 The first is a laterally oriented channel structure. In this device structure the photoactive material forms a channel, usually a few m wi de, between two electrodes. The quantum dot d evices published with this structure are photoconductors The biggest advantage to using this structure is that photoconductors using this stru cture benefit from high photoconductive gain, which in turn allows for high D* values. The gai n G is defined as the ratio of carrier lifetime to carrier transit time t as in The working principle is that upon photoexcitation, one carrier (usually the electron in PbS) is trapped in mid bandgap states [28] while the other carrier is extracted and allowed to be reinjected into the device through the circuit. Thus, whe n a device exhibits gain, the non trapped charge carriers effectively traverse the circuit m ultiple times

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46 The best photoconductive QD photodetectors [57,58] achieved the ~ 10 1 3 cmHz 1/2 /W level of InGaAs photodetectors However, the main disadvantage is that since the gain in quantum dot photodetectors is generated by the presence of traps, the devices using this stru cture have slow response times with bandwidths often limited to just tens of hertz [57,58] This limits practical applications to low frame rate imagi ng, static imaging, and other low speed applications. Another disadvantage of this structure is that the photoactive material is directly exposed to the environment unless there is some device encapsulation used. While recent synthet ic work has improved on the air stability of many QD systems, there are still concerns over chemical ( e.g. moisture or solvent exposure) and UV light induced degradation, which could play a role in this structure. The second device structure is the vertically stacked photodetect or. Devices with this structure can be photoconducting phot odetectors if there is gain, but are predominantly photodiodes without gain. In this structure, Schottky diodes [59] P N diodes [60,61] and P I N diodes [62] can be formed. The photodetectors in Chapter 5 of this dissertation are P I N like in structure. Here, the photoactive layer is sandwiched between two charge blocking layers that serve to reduce the dark current in the device. Lower ing the dark current by inserting charge blocking layers is advantageous because it lowers the noise floor and increases D* Also, if the charge blocking layers are air stable, they can serve as a native encapsulation for many different environmentally sen sitive photoactive layers. Early developments with the P I N structure utilized sensitive organic charge blocking layers, but found that by substituting appropriate metal oxides, not only was the dark current reduced even further, but the air stability of the device improved [62] Also, since these photodiodes have little to no photoconductive gain, they typically

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47 have faster operating speeds than photoconductors in the lateral channel configura tion. It is for these reasons why we chose to pursue the P I N structure in this dissertation. 1.5 Figures Figure 1 1. Three classes of organic molecules. Figure 1 2. Diagrams showing A) sp 2 hybridization with and bonding, and B) the energy level diagram for sp 2 hybridization [63]

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48 Figure 1 3. A diagram showing sp 2 hybridization and conjugation in a benzene ring Image extracted from wikicommons Figure 1 4. A diagram showing the three types of excitons in solid materials and their energetic distribution in the bandgap. Images extracted and modified from MIT Open Courseware lectures on Organic Electronics.

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49 Figure 1 5. An electron/polaron hopp ing between potential wells. Figure 1 6. The AM 1.5G solar spectrum.

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50 Figure 1 7. A generic current density voltage curve in the dark and under photoexcitation showing the short circuit current, open circuit voltage, and maximum power point. A B Figure 1 8. Device structures for A) the bilayer planar heterojunction solar cell and B) the modern bulk heterojunction solar cell.

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51 Figure 1 9 Synthesis setup for colloidal quantum dots. A three necked flask is loaded with one precursor solution under Ar flow with a thermocouple to track the temperature. The second precursor is injected, then extracted with a syringe [38]

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52 Figure 1 10 The chemical processes involved in the nu cleation and growth of colloidal quantum dots in the hot colloidal injection technique [38] Figure 1 11 Free energy considerations in the nucleation phase of colloidal quantum d ot synthesis. The plot shows the interplay between the surface free energy cost and volume free energy gain by forming nuclei [49]

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53 Figure 1 12 The relationship between growth rate and size for different precursor solution concentrations. The critical size changes with precursor solution concentration [47] Figure 1 13 The energy level evolution as atoms are assembled from one single cation anion pair t o the bulk band structure. The band structure of quantum dots lies between the molecular orbitals and bulk band structure.

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54 Figure 1 14 Tunability of the quantum dot bandgap is shown in different batches of PbSe synthesized for quantum dot based light emitting diodes. The absorption edge changes as a function of quantum dot size [45] Figure 1 15 A schematic of an electron hopping between discreet energy states. This simulates the hopping transport mechanism active in quantum dot transport.

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55 Figure 1 16 This plot shows the working wavelengths for all the common infrared photo detector materials. MCT = HgCdTe. Figure extracted from Teledyne Judson Technologies website [64] Figure 1 17 A schematic of the complex structure of expitaxia lly grown infrared detectors in a focal plane array mounted to the silicon read ou t circuit by indium bump bonding. The Si ROIC is typically a CMOS circuit, involving many processing steps by itself.

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56 Figure 1 18 A comparison of the lateral channel stru cture and vertical stack device structure for photodetectors. The structure on the left is the lateral channel structure common in photoconductors. The structure on the right is the vertical stack structure common in photodiodes without gain.

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57 CHAPTER 2 SOLUTION PROCESSED NICKEL OXIDE HOLE TRANSPORT LAYERS IN HIGH EFFICIENCY PHOTOVOLTAIC CELLS 2 1 Background In a society concerned with environmental, economic, and geopolitical consequences of energy consumption, new alternatives to traditional energy sources are needed. With rapid progress being made in organic photovoltaics they are becoming a viable source of renewable energy as power conversion efficiencies (PCEs) exceeding 8% have been demonstrated [6,16] Traditional bulk heterojunction polymer solar cells con sist of a transparent indium tin oxide (IT O) anode a hole transport layer, a photoactive layer, and a top cathode. H ole transpor t layers must have a high optical transparency good chemical stability, a large ionization potential and good electron blocking capability In a typical polymer solar cell, PEDOT:PSS is used as the HTL and has a work function ( of 5.2 eV. However, its acidity, tendency to absorb water, and inability to block electrons effectively are factors which contribute to device p erformanc e problems and degradation [65] Nickel oxide is emerging as an alternative HTL for polymer solar cells [25,26,66 74] Pure, s toichiometric NiO is an excellent insulator, with room temperature conductivity on the order of 10 13 S cm 1 [75] whil e non stoichiometric NiO is a wide bandgap p type semiconductor [76 81] The p type conductivity of NiO originates from two positively charged holes which accompany each Ni 2+ vacancy in the lattice for charge neutrality [76,82,83] These holes ar e quasi localized on Ni 2+ ions near the vacancy in the lattice, generating two Ni 3+ ions for each Ni 2+ vacancy [76,84] The valence band edge of NiO is well a ligned to the highest occupied molecular orbital (HOMO) levels of many p type conjugated polymers for photovoltaics [74,81] Irwin et al. first demonstrated an enhancement in p olymer solar cell performance with a NiO electron blocking layer deposited via pulsed laser deposition [74,85] Recently, solution

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58 processed NiO was also reported for polymer photovoltaics [25,26,81] A nickel ink made from nickel formate and ethylenediamine was used as the precursor in those reports. Here, we chose nickel acetate tetrahydrate and monoethanolamine precursors in an ethanolic solution, as this presents a set of materials not yet used to create NiO in solar cells. Using these NiO solution precursors, we investigated the chemical, optical, electronic, and structural properties of solution processed NiO from conception of the precursors to their impact on the solar cell. Combining this solution derived NiO with our low bandgap polymer poly dithienogermole thienopyrrolodione ( p DTG TPD), [6,23] we fabricated solar cells with an average PCE of 7. 8 %. Th is is a 1 5% enhancement over reference devices with PEDOT:PSS as the HTL. Contrary to previous reports, [25,26,66,67,74,81] we found that neither the effective work function difference between NiO and PEDOT:P SS nor the manifestation of the better electron blocking ability of NiO in the reverse saturation current were the most important factors in the device improvement. Rather this solution derived NiO provided a more favorable surface to form optimized donor /acceptor nanoscale phase morphology at the HTL/BHJ interface, reducing the series resistance and increasing the shunt resistance in the device. Additionally, optical absorption in the solar cells is improved due to an optical resonance shift when NiO is s ubstituted for PEDOT:PSS. Finally the devices with NiO were found to be more air stable than devices with PEDOT:PSS under ambient conditions. 2.2 Results and Discussion 2.2.1 NiO Precursor Composition The precursors used to synthesize common solution de rived metal oxides consist of a solution of a metal salt and an amine compound in an alcohol solvent. The amine compounds for m etal oxide synthesis are often considered to be a sol modifier or stabilizer [86 88] In contrast to these sugg estions, we found that the amine plays an active role in the precursor solution, forming

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59 ionic complexes with the metal, in agreement with previous studies [89] rather than a passive stabilizer role. To synthesize the precursor solution, we dissolved nickel acetate tetrahydrate (Ni(OAc) 2 4H 2 O) and monoethanolamine (MEA) in ethanol. To determine the exact role MEA plays in the reaction, high resolution electrospray ionization mass spectrometry (ESI MS ) was used to sample the ionic species in the solution The main species found in the precursor solution was the [Ni( MEA ) 2 ( OAc)] + ion which has a mass to charge ratio (m/z) of 239.0520. The composition of this ion was confirmed by comparing the experimenta l data to the theoretical mass to charge ratio as shown in Figure 2 1 and Figure 2 2. The ionic structure we concluded is based on octahedral Ni bonding and the UV Vis NIR absorption spectrum, which has an absorption minimum in the green region and maximum in the red and violet regions of the visible spectrum. According to Ligand Field Theory, nickel complexes composed of oxygen donor atoms appear green, while those with nitrogen donor atoms appear from deep blue to pink, depending on the coordination numbe r and type of amine. Since the solution appears deep green and the absorption spectrum shows the absorption pattern of octahedral Ni species, we conclude the coordination must be bidentate for all ligand species, with oxygen composing four out of six atoms coordinated to nickel. This bidentate coordination creates three low energy ring structures and pushes the visible absorption to red and violet, creating a deep green solution. From these data, we conclude that MEA acts as a ligand that coordinates to ni ckel to form ionic complexes and plays a larger role than a solution stabilizer. Therefore, the formation of NiO requires complete thermolysis of the ionic complexes, the details of which will be presented in the following section.

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60 2.2.2 Characterization of NiO Films 2.2.2.1 Optical p roperties Figure 2 3 shows the UV Vis NIR absorption spectra of the precursor solution and the NiO films fabricated at different temperatures. Nickel oxide films were formed by thermally decomposing spin cast films of the pre cursor solution in air. The absorption spectrum of the precursor solution is consistent with the typical d d transitions of octahedral nickel complexes coordinated mostly by oxygen donor atoms in agreement with the chemical structure deduced from mass spe ctrometry When spin cast precursor films were heated to 230 o C, the absorption corresponding to the precursor remained in the spectrum while a strong absorption beginning at 330 nm emerged, corresponding to the bandgap edge of NiO. These results indicate that the precursor is not completely converted to NiO at 230 o C. When the films were heated to 275 o C, the precursor was completely converted into NiO, as shown by the strong band edge absorption at 330 nm and the absence of precursor absorption The broad absorption tail at energies below the bandgap for NiO fabricated at 275 o C is caused by nickel vacancies, resulting in a slight colora tion in the films [90,91] This coloration in th in films does not significantly hinder the optical transparency, which is important for solar cell performance. Figure 2 4 shows the optical transmission spectrum of a 5 nm thick solution derived NiO film a long with the spectrum for a 30 nm thick PEDOT:PSS film used as a reference These were the thicknesses of the HTL films that gave optimized performance in the solar cells discussed later. The NiO film is more than 95% transparent across the entire visible and near infrared spectra. At wavelengths longer than 590 nm, where the solar flux is largest, the NiO film is more transparent than the PEDOT:PSS film which is desirable for solar cell applications. These results suggest that NiO is a promising candidate for HTL in OPV cells.

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61 2.2.2.2 Electronic p roperties To understand the composition of the solution derived NiO films fabricated at 275 o C, these films were studied by X ray photoelectron spectroscopy (XPS). Figure 2 5 shows the XPS spectrum for the Ni 2p 3 /2 state w hich can be separated into four peaks. First, the peak centered at a binding energy of 853.7 eV corresponds to Ni 2+ in the standard Ni O octahedral bonding configuration in cubic rocksalt NiO [92 94] Second, the broad peak centered at 860. 7 eV has been ascribed to a shake up process in the NiO structure [93] The peak centered at 855.3 eV has been ascribed to the Ni 2+ vacancy induced Ni 3+ ion [92,93,95] or nic kel hydroxides [93,96 100] The peak centered at 856.5 eV signifies nickel oxyhydroxide (NiOOH), a critical surface dipolar species in high performance polymer solar cells [26,100] Figure 2 6 shows the XPS spectrum for the O 1s state, which can be separated into two distinct peaks. The peak centered at 529.2 eV confirms the Ni O octahedral bonding in NiO [93] The peak at 530 8 eV is ind icative of nickel hydroxides including defective nickel oxide with hydroxyl groups adsorbed on the surface [99 101] However, this peak has also been assigned to oxygen interactions with a nickel deficient lattice, suggesting some correlation wi th nickel vacancy concentration [93,95] The peak at 531.7 eV is the partner to the 856.5 eV peak, representing NiOOH [100] Since there are several possible sources of the XPS signals at 855.3 eV and 530.8 eV we further characterized the structural and chemical features of the film by grazing incidence X ray diffraction (GIXRD) and angle resolved XPS to examine which of these sources are present, if not all of them. Detailed analysis of the NiO films by GIXRD shown in Figure 2 7 showed the typical NiO d iffraction pattern with small grains approximately 1 nm in diameter, as estimated from the Scherrer equation [102] T he presence of nickel vacancies is consistent with shifted diffraction peaks relative to those of stoichiometric NiO Compa red to the simulated bulk

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62 stoichiometric NiO, there is a shift to lower 2 (larger d spacing) for all of the experimentally determined diffraction peaks. This is caused by the presence lattice strain produced by electrostatic repulsion in the lattice. This electrostatic repulsion is caused by the presence of nickel vacancies and/or oxygen interstitial s. This is further explained as the following: the nickel vacancies reduce the Coulombic screening between oxygen ions in the lattice, so the oxygen ions repel each other, increasing the d spacing. In the case of oxygen interstitials, the strain is induced by the lattice accommodating the additional atom in an interstitial site and its additional Coulomb interactions. This effect has been shown in NiO crystal gr owth using other methods by varying the oxygen partial pressure in the growth ambient and measuring the Ni:O stoichiometry [103,104] Peak shifts due to X ray refraction at small incident angles have als o been taken into consideration [105] Angle resolved XPS spectra of the O 1s and Ni 2p 3/2 regions in Figure 2 8 show an incr eased i ntensity of the 855.3 eV and 530.8 eV peaks at a lower take off angle due to the presence of hydroxides a t the extreme surface lay ers caused by air exposure The NiOOH peak also increases at low take off angles, confirming the surface presence. This species, in particular, was found to be an essential surface dipolar feature for favorable contacts in bulk heterojunction PV cells [100] Additionally, as discussed above, the NiO films do have a pale black tinge, which further confirms the vacancy induced Ni 3+ and NiOOH [90,91,106] 2.2.3 Photovoltaic Cells with Solution Processed Nickel Oxide Using the NiO films described above as a hole transp orting layer, we fabricated bulk heterojunction solar cells based on the high efficiency pDTG TPD:PC 71 BM photoactive layer. The device structure and the energy band diagram of the NiO device are shown in Figure 2 9 along with that of the reference PEDOT:PS S device [23,107,108] The device fabrication details are reported in the experimental procedure section.

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63 2.2.3.1 Solar cell c haracteristics To demonstrate the viability of these NiO films for photovoltaic applications, we fabricated solar cells with NiO as the HTL with three different NiO precursor processing temperatures: below the completed precu rsor thermolysis at 185 o C, in the middle of thermolysis at 230 o C, and after complete thermolysis of the NiO precursor films at 275 o C. The current density voltage ( J V ) characteristics of these solar cells are presented in Figure 2 10 with the performanc e characteristics summarized in Table 2 1 At a processing temperature of 185 o C, the precursor did not significantly decompose into NiO, resulting in a non conductive HTL and poor solar cell performance, as expected. At 230 o C, the HTL film contained a mixture of residual non conductive precursor and NiO, and thus the solar cell showed an improvement in short circuit current density ( J sc ) and open circuit voltage ( V oc ). When heated to 275 o C, the film fully converted to NiO and the solar cell showed opti mum performance with an average PCE of 7.8%. We note that in contrast to previous reports of solution processed NiO in PV cells, where oxygen plasma treatment of NiO is needed for optimized devices [25,26,81] the NiO films used here did not require any post de position surface treatment for optimized performance I n fact, all of the solar cell performance parameters decreased when UV O 3 treatment was applied to NiO before depositing the active layer as shown in Figure 2 11 We benchmarked the performance of o ur NiO based devices by fabricating them side by side with PEDOT:PSS based devices. The J V characteristics of the solar cells with either an optimized 5 nm thin NiO HTL or an optimized reference 30 nm thin PEDOT :PSS film are shown in Figure 2 12 The perf ormance characteristics are summarized in Table 2 2 The open circuit voltages of the devices with PEDOT:PSS were 10 mV higher than those with NiO. Devices with PEDOT:PSS produced an average J sc of 12.7 mA cm 2 while those with NiO produced an average of 13.9 mA cm 2 a 9.4 2.9 % improvement. The average fill factor ( FF ) for devices

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64 with PEDOT:PSS was 64.2%, while for those with NiO it increased to 68.4%, a 6.5 0.8 % improvement. Because of the enhanced J sc and FF devices with NiO showed significantly higher PCEs than those with PEDOT:PSS. The optimized PEDOT:PSS based device showed an average PCE of 6.8% while the optimized NiO devices showed an average PCE of 7.8%, a 14.7 0.1 % improvement. To understand why the solar cells performed as such relative to each other, we studied their performance parameters in more detail. 2.2.3.2 Analysis of s olar c ell p erformance In previous work with solution processed NiO, the V oc of the NiO devices was higher than that of the PEDOT:PSS devices due to the larger work function of NiO and increase in electron blocking capability, i.e. lower dark reverse saturation current ( J o ). [26,81,100] In the limit of large shunt resistance (which is applicable here and is discussed later), the relationship between V oc and J o described in the ideal diode model [81] can be written as : (2 1) where n is the diode ideality factor, k T is the absolute temperature, and q is the magnitude of the elementary charge. Our device data showed that devices with NiO have a lower J o than the devices with PEDOT:PSS. This suggests that there is improved electron blocking in the NiO devices. Based on Eq uation 2 1 a lower J o s hould also lead to a higher V oc which is not the case here. Furthermore, the magnitude of the predicted V oc in these devices does not agree with this relation when we insert the measured values for the J sc J o and n into Eq uation 2 1 Fitting the dark current data yields a diode ideality factor in both devices of 1.27 as shown in Figure 2 13 As an example, we calculated the expected V oc using the J o values at 0.5 V: J o for PEDOT:PSS devices is 2.5x10 5 mA/cm 2 while for N iO devices it is 1.7x10 5 mA/cm 2 The result is a calculated V oc for the PEDOT:PSS devices of 0.44 V, while for the NiO

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65 devices it is slightly higher at 0.46 V. Clearly, this is not in agreement with the experimental data, either in magnitude, or in that P EDOT:PSS based devices have a larger V oc experimentally. Furthermore, our results indicate that the ideal diode equation is unable to fully account for the device behavior, suggesting the behavior of the solar cells is more complicated than this simple mo del. To fully explain the difference in V oc and ascertain the source of the enhanced short circuit current density and fill factor in our devices, we measured the built in voltage ( V bi ) of the solar cells with electroabsorption (EA), studied the spectral response of the solar cells and the morphology of the active layer at the HTL/active layer interface. The built in voltage of solar cells is the difference in the effective work functi ons of the cathode and anode/ HTL. Using EA, we determined that the buil t in voltage in the NiO and PEDOT:PSS devices was 0.95 V, and 0.96 V resp ectively as shown in Figure 2 14 The built in voltage of the solar cells is determined by extrapolating the linear regime of the to the horizontal axis at T/T = 0. Thus, if we tak e the work function of PEDOT:PSS to be 5.20 eV, as stated above, then the effective work function of the NiO in the device is 5.19 eV, as extracted from the EA data. T his 10 mV difference in the V bi is the same as the 10 mV difference in the V oc of the two solar cells. It is reasonable to conclude that changing the work function of the HTL plays a role in controlling the V oc in this system since the HOMO level of pDTG TPD lies below the Fermi level of NiO that has been fabricated outside of vacuum condition s [23,24] Similar modulation of the V oc by modification of the HTL work function was realized with PCDTBT based solar cells, where PCD TBT also has a deep HOMO level [81] However, in cases where P3HT was used as the donor polymer in the BHJ, no f urther modulation of the V oc was observed when increasing the work function of the HTL. This was because the HOMO level of P3HT lay closer to the vacuum level than the Fermi level of the HTL [25,81,109] That is, modulating the V oc

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66 by changing HTL work function seemed to occur only when the HOMO of the donor polymer lies below the Fermi level of the HTL. Based on these results, we conclude that the V oc is controlled by the work function of the HTL in this system since the Fermi level of both HTLs lies above the HOMO level of the donor polymer, and that the enhancements in J sc and fill factor in the NiO cells are not due to the difference in work functio n of the HTLs. To determine the cause of increased J sc we examined the external quantum efficiency (EQE) and optical resonance in the devices. The EQEs for both devices are shown in Figure 2 15 The global maximum EQE is in the NiO device and has a valu e of 67% at 615 nm, which is in agreement with the slightly higher optical transmission of NiO in this wavelength range. The maximum EQE for PEDOT:PSS based devices is slightly lower, with a value of 65% at 610 nm. Interestingly, w hile the transmission of NiO is slightly lower than that of PEDOT:PSS at wavelength s between 425 590 nm, the EQE in th e NiO devices is actually higher. This enhancement in the EQE is due to the different refractive indices of the HTLs and o ptical resonance in the devices and is one of the main reasons for the enhancement in J sc in the NiO devices. The optical resonance change has been confirmed by our Finite Difference Time Domain (FDTD) method optical modelin g results, as shown in Figure 2 16 J o are not the primary reasons for the enhancement in overall device PCE, and the fill factor improvement is not directly correlated with the optical resonance shift, we stu died the differences in surface morphology and surface energy between NiO and PEDOT:PSS. Based on the atomic force microscopy (AFM) results, the root mean square (RMS) roughness values for PEDOT:PSS and NiO were found to be 0.79 nm and 0.88 nm, respectivel y. This small difference in roughness

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67 should not affect the solar cell performance. In previous work by Hau et al ., changes in the morphology of the active layer in inverted solar cells were correlated with changes of surface energy of the TiO 2 or ZnO elec tron transport layers (ETLs) on which the active layer was deposited [110,111] Depositing various self assembled monolayers on the ETLs prior to depositing the active layer modified the surface energy of t he ETLs. These active layer morphology changes yielded an increase in short circuit current and fill factor in the solar cells [110,111] This type of analysis has not been per formed to explain the performance of solar cells with NiO as a HTL. We examined the surface energy of the HTLs by measuring water contact angles of the NiO and PEDOT:PSS films. The avera ge contact angles were 29.3 2.8 for NiO and 12.5 1.4 for PEDOT:P SS as shown in Figure 2 17 These results indicate that the NiO surface is less hydrophilic than PEDOT:PSS, allowing better wetting by nonpolar solvents such as that used for the pDTG TPD/PC 71 BM active layer. This improved wetting leads to improved active layer film formation and possibly donor/acceptor phase morphology, as observed in previous work described above. To confirm the improvement in donor/acceptor phase morphology near the HTL/active layer interface, we deposited a 10 nm thick active layer fil m on both PEDOT:PSS and NiO surfaces Figure s 2 18 a,b show the atomic force microscopy (AFM) roughness results. The RMS roughness of the thin active layer film on NiO was 2.3 nm, while the roughness of the film on PEDOT:PSS was 3.6 nm. This indicates that the film formation and physical contact at the interface of the active layer and the HTL is more uniform when using NiO. To further examine the chemical differences at the interface, we examined the phase distribution of these thin BHJ films. As shown in Figures 2 18 c,d the active layer film deposited on NiO has a more ho mogeneous phase distribution near HTL/active layer interface than the film deposited on PEDOT:PSS. This overall improved morphology and physical contact with the HTL s hould lead

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68 to a more uniform electrical contact to NiO. In fact, there is a 40 % lower series resistance ( R s ) and 31 % higher shunt resistance ( R sh ) in the NiO devices compared to the PEDOT:PSS devices as shown in Table 2 2. The decrease in series resistance indicates that holes are extracted more efficiently, while t he increase in shunt resistance suggests that there is less interfacial recombination in the NiO device and that NiO is blocking electrons more effectively It is possible that both the energy band alignment of NiO with the BHJ HOMO LUMO levels and the improved donor/acceptor phase morphology are both contributing to the change in parasitic resistances in these devices. In this case, the NiO is actually mitigating the parasitic resistances and i mproving fill factor due to improved donor/acceptor morphology at the NiO/BHJ interface and improved contact. This is where the electron blocking ability of NiO is expected to have its most pronounced effect. Additionally, we found that there is a change in the morphology as a function of distance from the HTL/BHJ interface. An inspection of the ful l ~100 nm thick BHJ layer by AFM revealed nearly identical RMS roughness values of approximately 3.1 nm for the film on NiO and 3.2 nm for the film on PEDOT:PS S and nearly identical phase mapping for both BHJs deposited on PEDOT:PSS or NiO, as shown in Figure 2 19 This suggests that the donor/acceptor morphology is strongly affected by the HTL surface energy at the HTL /BHJ interface while it is weakly, if at a ll, affected far away from the interface A detailed study of the vertical donor/acceptor phase separation is beyond the scope of this work, but would help elucidate the exact interfacial science present here. Here, we conclude that this improved donor/acc eptor phase morphology at the HTL/BHJ interface and enhanced optical absorption in the device are the major factors in the success of the NiO device.

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69 2.2.3.3 Stability of solar cells The final indicator of the viability of NiO in solar cells is the devic e storage stability. We performed a side by side comparison of un encapsulated NiO and PEDOT:PSS based devices Figure 2 20 the NiO based devices degraded much slower than the PEDOT:PSS devices. T his is in agreement with previous work on degradation of solar cells with NiO compared to PEDOT:PSS [67,81] Most of the difference in performance loss between the two devices was due to the decrease in J sc indicating the difference in device degradation is due to HTL and/or interfacial degradation Indeed, previous studies have shown that the acidity and hygroscopic nature of PEDOT:PSS corrodes the ITO electrode and PEDOT:PSS itself [65] ; both are problems for organic photovoltaics at this time which can be overcome by substituting a more stable mat erial for PEDOT:PSS. Thus, we attribute the quicker degradation of the PEDOT:PSS based device to the more acidic and hygroscopic nature of PEDOT:PSS compared to NiO. 2.3 Conclusions We characterized the physical, chemical, optical, and electronic propertie s solution processed NiO and fabricated solar cells using this NiO film as a ho le transport layer. We found that ethanol soluble monoethanolamine coordination complexes of nickel can be thermally decomposed to form NiO. XPS results showed a strong contribu tion from the Ni 3+ state, confirming the formation of p type, non stoichiometric NiO and the dipolar NiOOH species upon thermolysis of the precursor films. Optical analysis of the NiO thin films showed favorable optical properties for photovoltaic applicat ions. Solar cells incorporating NiO films showed significant enhancements in fill factor and short circuit current, leading to a 14.7% increase in PCE compared with the cells with PEDOT:PSS. The enhancements in the NiO devices were due

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70 to improved optical resonance, nanoscale active layer morphology, increased shunt resistance, and lower series resistance for charge extraction in the NiO devices. 2.4 Experimental Procedure NiO precursor solution. Nickel acetate tetrahydrate (Ni(CH 3 COO) 2 4H 2 O) (Acros Organ ics) was dissolved in ethanol with monoethanolamine (NH 2 CH 2 CH 2 OH) (Sigma Aldrich) (0.1 mol L 1 ). The mole ratio of Ni 2+ : MEA was maintained at 1:1 in solution. Dissolution took place while stirring in a sealed glass vial under air at 70 C for four hours. The solution appeared homogeneous and deep green after approximately forty minutes. NiO film preparation All NiO films were deposited by spin casting onto the appropriate substrates. For solar cells, the substrates were ITO coated glass with sheet resista nce of 15 ohm sq 1 For GIXRD, the substrates were plain glass. For XPS, the substrates were single crystal Si wafers. All substrates were cleaned by the following procedure before spincasting: successive sonication in DI water, acetone, and isopropanol ba ths for 15 minutes each. Unless otherwise stated, the NiO precursor films were heated to 275 C for 45 minutes in air. Device f abrication and c haracterization To fabricate the solar cells, ITO coated glass substrates were UV ozone treated for 15 minutes between solvent cleaning and spin casting of the HTL. We noted, however, that this UV ozone treatment of ITO was not necessary for proper wetting by the NiO pr ecursor solution. PEDOT:PSS (Clevios Al 4083 ) was filtered through a 0.45 m PTFE filter prior to spin casting at 7000rpm. The PEDOT:PSS films were heated at 140 o C for 10 minutes in air and transferred immediately to the nitrogen filled glovebox for furth er processing. When NiO was used as the HTL instead of PEDOT:PSS, the NiO films were allowed to cool before being transferred to a nitrogen filled glovebox for further processing. The polymer:fullerene blend ratio is 1:1.5 by weight and is dissolved in chl orobenzene with 5 vol%

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71 1,8 diiodooctane (DIO) additive. PC 71 BM was acquired from Solenne (purity >99%) and was used as received. The blend was spin cast on top of the HTL film to give an active layer thickness of 105 nm. The devices were immediately loaded into a thermal evaporator where 1 nm of LiF and 100 nm of Al were deposited under a vacuum pressure of 5x10 7 Torr. Current voltage ( J V ) characterization was performed with a Keithley 4200 semiconductor parameter analyzer system with a Newport Thermal Or iel 94021 1000 W solar simulator ( 4 in. by 4 in. beam size) using the AM1.5 G solar spectrum at 100 mW cm 2 incident power The light intensity was calibrated by a n ORIEL 91150V monosilicon reference cell calibrated by Newport Corporation. EQE measurement s w ere conducted using a n in house setup consisting of a Xenon DC arc lamp, a n ORIEL 74125 monochromator, a Keithley 428 current amplifier, a n SR 540 chopper system and a n SR 830 DSP lock in amplifier from SRS The experimental setup for electroabsorption was described previously. [112] GIXRD measurements GIXRD measurements were performed on a Philips XPert MRD with Cu K X rays (non monochromatic) with incident angle = 3. XPS measurements XPS data were acquired on a Perkin Elmer 5100 XPS system with a non monochromatic Al anode X ray source with a pass energy of 35.75 eV Peak deconvolution and analysis was performed in RBD AugerScan version 3.2 software. The adventitious C 1s peak was referenced to 284.8 eV. Mass s pectrometry measurements Solutions were introduced into an Agilent 6210 TOF MS via direct injection followed by electrospray ionizat ion (ESI) with an autosampler. The mobile phase was spectroscopy grade ethanol from Fisher Scientific and was used as received. Accurate mass identification was performed in MassHunter software.

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72 2.5 Figures and Tables Figure 2 1. High resolution electrospray ionization mass spectrum of the precursor compound. The peak family shown corresponds to the isotopic distribution of nickel; the ion at m/z = 239.0520 contains the most common isotope, Ni 58. The inset is the proposed structure of the ion. Figure 2 2. Theoretical high resolution mass spectrum for [Ni(MEA) 2 (OAc)] + This spectrum is in agreement with the experimentally acquired pattern for the [Ni(MEA) 2 (OAc)] + ion, confirming the unique mass assignment.

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73 Figure 2 3. Absorption spectra of the transformation of the precursor into NiO. The material transforms from the precursor to a full nickel oxide film by thermal decomposition of the precursor materials, as is shown by the elimination of the visible and IR absorption peaks, and the appearance of the UV band edge of NiO. Figure 2 4. Transmission spectrum of 5 nm thin NiO films is compared to 30 nm thin PEDOT:PSS films used in the optimized solar cells. The thin NiO films are more transparent than the PEDOT:PSS films used in solar cells at wav elengths longer than 590 nm.

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74 Figure 2 5. High resolution Ni 2p 3/2 XPS acquisition f or NiO. The spectrum shows four contributions one from Ni 2 + in the octahedral NiO configuration at low binding energy (red) one from hydroxylated or defective NiO at an intermediate binding energy (blue) one from nickel oxyhydroxide (blueish green) and a high energy peak from a shake up process (light green) in the NiO lattice at the highest binding energy. Figure 2 6. High resolution O 1s XPS acquisition for NiO. Th e spectrum shows three contributions one from O 2 in the octahedral NiO configuration at low binding energy (red) and one from hydroxylated or defective NiO (blue) and one from an oxyhydroxide at a higher binding energy (blueish green)

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75 Figure 2 7. GIXRD spectrum of solution derived NiO, heated to 275 o C for 45 minutes on a glass substrate. The bulk stoichiometric NiO diffraction pattern is also shown for comparison. The shift of the experimentally acquired GIXRD spectrum is indicative of crystallographic vacancies and interstitials. Figure 2 8. Angle resolved XPS spectra for the Ni 2p 3/2 and O 1s states. At low take off angles, the peak s corresponding to hydroxylated (blue) and oxyhydroxylated (bright green) NiO increase in intensity r elative to the NiO peak (dark red) corresponding to extreme surface positions for the hydroxylation.

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76 Figure 2 9. Device structure and energy levels with respect to vacuum of the materials used in the photovoltaic cells. Figure 2 10. Illuminated curren t density voltage ( J V ) characteristics of the solar cells with NiO fabricated at different temperatures. After the full formation of NiO at 275 o C the solar cells perform optimally, with a PCE of 7.8%.

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77 Figure 2 11. A) J V characteristics of solar cells fabricated with varying UV O 3 treatment time on NiO HTLs. Any UV O 3 treatment on NiO decreased the device performance compared to devices without any treatment. B) High binding energy cut off in the XPS survey scan of the freshly prepared NiO f ilm with that of a NiO film after 10 minutes of UV O 3 treatment on freshly prepared single crystal Si substrates cleaved from the same wafer. The data show that there is a work function shift of more than 1 eV with UV O 3 treatment. This is cannot be direct ly correlated to the ~0.3 V change in Voc of the solar cells Data were corrected for sample charging by referencing the adventitious C 1s peak to 284.8 eV Figure 2 12. Illuminated current density voltage ( J V ) characteristics of solar cells comprising NiO or PEDOT:PSS HTLs. The solar cells with NiO as the HTL outperform those with PEDOT:PSS due to a higher fill factor and short circuit current. A B

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78 Figure 2 1 3 The plot of diode ideality factor vs. voltage for both types of solar cells. Figure 2 14. Built in voltage of solar cells with either NiO or PEDOT:PSS as the HTL determined by electroabsorption. The built in voltage of the solar cells differs by the same 10 mV by which the V oc differs, suggesting the V oc is limited by the built in voltage in these solar cells.

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79 Figure 2 15. External quantum efficiency of solar cells fabricated in this study. The EQE of solar cells with NiO is greater at nearly all wavelengths of incident light. The spectrum shift between the device is caused by opt ical resonance changes within the solar cell. Figure 2 16. Simulated spectral absorption for solar cells with NiO or PEDOT:PSS HTLs. The spectra are in close agreement with the experimental spectra, indicating that the shift is due to optical resonance i n the solar cells.

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80 Figure 2 17. Water contact angles for NiO and PEDOT:PSS HTLs. The avera ge contact angles were 29.3 2.8 for NiO and 12.5 1.4 for PEDOT:PSS The NiO HTL is less hydrophilic, promoting favorable wetting during solution processing and BHJ physical contact. Figure 2 18 AFM roughness images of 10nm thick BHJ films on a) NiO and b) PEDOT:PSS showing the contrast in physical formation of the films on different HTLs. The active layer films deposited on NiO are smoother than those depo sited on PEDOT:PSS. AFM phase images of thin polymer/fullerene blend layers on c) NiO and d) PEDOT:PSS, showing a drastic change in material distribution near the HTL/active layer interface, caused by surface energy differences of NiO and PEDOT:PSS. The sc ale bar is 400nm and the scan area is 2 m x 2 m. [113]

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81 Figure 2 19 AFM roughness images of the full ~ 100 nm thick BHJ films on a) NiO and b) PEDOT:PSS AFM phase images of the full ~100 nm thick BHJ films on c) NiO and d) PEDOT:PSS. The images show nearly identical roughness and phase mapping. The scale bar is 400nm and the scan area is 2 m x 2 m.

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82 Figure 2 20 Normalized device characteristi cs as a function of air exposure time for (a) PEDOT:PSS and (b) NiO based devices. Devices made with NiO HTLs are more stable when stored in ambient conditions, indicating that the NiO HTL is more stable than the PEDOT:PSS HTL in these solar cells. Table 2 1. Performance of solar cells fabricated at various NiO processing temperatures. One standard deviation is reported in parenthesis. NiO Temperature V oc (V) J sc (mA cm 2 ) FF (%) PCE (%) 185 o C 0.60 0.09 (0 .01 ) 13.0 ( 0.1 ) 0.007 (0 .001 ) 230 o C 0.76 6.4 ( 1.9 ) 12.1 ( 0.8 ) 0.61 (0 .2 ) 275 o C 0.82 13.9 (0 .3 ) 68.4 ( 0.4 ) 7.82 (0 .2 ) Table 2 2. Device characteristics of the solar cells fabricated in this study. One standard deviation is reported in parenthesis. Series resistance was calculated using the illuminated J V values converging to V oc Shunt resistance was calculated using il luminated J V values converging to J sc HTL V oc (V) J sc (mA cm 2 ) FF (%) PCE (%) R s ( cm 2 ) R sh ( cm 2 ) V bi (EA) PEDOT:PSS 0.83 (0.005) 12.7 (0.2) 64.2 (0.3) 6.8 (0.1) 13.1 (1.1) 348.4 (90.4) 0.96 V 5 nm NiO 0.82 (0.006) 13.9 (0.3) 68.4 (0.4) 7.8 (0.2) 7.9 (1.3) 455.4 (72.5) 0.95 V

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83 CHAPTER 3 NOVEL LOW TEMPERATURE ROUTE TO SOLUTION PROCESSED NICKEL OXIDE HOLE TRANSPORT LAYERS IN POLYMER PHOTOVOLTAICS 3.1 Background Interest in new renewable energy sources is spurring growth in organic photovoltaics. The performance of organic photovoltaic cells is quickly improving, as power conversion efficiencies (PCEs) exceeding 8% have been demonstrated in single cells [6,16] and over 10% in tandem cells [114] T ypical bulk heterojunction polymer solar cells consist of a transparent indium tin oxide (IT O) anode a hole transport layer, a photoactive bulk heterojunction layer, and a top reflective cathode. H ole transpor t layers in high efficiency solar cells must be optically transparent, chemically stable, have a large ionization potential and good electron blocking capability. In a typical polymer solar cell, poly(3,4 ethylenedioxythiophene) doped with poly(styrenesulfonate) ( PEDOT:PSS ) is used as the HTL a nd has a work function ( of 5.2 eV. However, PEDOT:PSS is known to etch the ITO anode on which it is deposited it is highly hygroscopic and does not effectively block electrons [65] These properties all contribute to device performance problems and degradation [65] An emerging replacement HTL in polymer solar cells is n ickel oxide [25,26,66 74] S toichiometric NiO is an insulator, with room temperature conductivity on the order of 10 13 S cm 1 [75] while non stoichiometric NiO is a wide bandgap p type semiconductor [76 81] The p type conductivi ty of NiO is generated from two positively charged holes which accompany each Ni 2+ vacancy in the lattice for charge neutrality [76,82,83] These vacancy induced holes ar e quasi localized on Ni 2+ ions near the vacant Ni latt ice site generating two Ni 3+ ions for each Ni 2+ vacancy [76,84] The valence band edge of NiO is well aligned to the highest occupied molecular orbital (HOMO) levels of many modern p type conjugated polymers for photovoltaics [22,74,81]

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84 Irwin et al. first demonstrated an enhancement in polymer solar cell performance with a NiO electron blocking layer deposited via pulsed laser deposition [74,85] publication of NiO in solar cells, several groups have used NiO deposited from a variety of techniques, including sputtering [67,68,71,72] PLD [66] and O 2 plasma treatment of evaporated Ni films [73] Every report which compared devices with NiO to a standard showed either comparable, or an improvement in, device pe rformance and temporal stability upon the incorporation of NiO While NiO films can easily be fabricated by a wide variety of these techniques, the future of organic photovoltaics relies on the ability of materials used to be compatible with large scale, h igh throughput techniques; thus the solution route to NiO is preferred and must be explored. Rec ently, solution processed NiO HTLs have made an appearance in polymer photovoltaic cells There have been several routes to employing solution processed NiO in solar [25,26,81,115] peptized nickel hydroxide in a standard device structure [70] and a dispersion of fine nickel oxide powder into a solvent in an inverted solar cell [69] The most common and reliable has been the sol gel route, in which a nickel salt is added in an alcohol solvent with an amine compound to form soluble coordination complexes NiO which need high tempera tures to decompose to yield NiO [25,26,81,115] In all of the studies of solution processed NiO with high efficiency solar cells, the researchers have employed high temperature processing steps to form NiO from the precursor solutions that are not compatible with polyethylene terephthalate (PET) substrat es envisioned for large scale roll to roll (R2R) processing methods. In chapter 2 we reported high efficiency polymer solar cells with solution processed NiO HTLs based on a BHJ blend of poly dithienogermole thienopyrrolodione ( p DTG TPD) and the acceptor fullerene [6,6] phenyl C 71

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85 butyric acid methyl ester (PC 71 BM) [115] In that study, a precursor solution of nickel acetate tetrahyd rate and monoethanolamine contained [Ni(MEA) 2 (OAc)] + ions, which thermally decomposed and oxidized at temperatures too high for PET substrates. To fabricate the NiO HTL at temperatures compatible with PET substrates, we hypothesized that a two step NiO for mation process was possible: 1) Preliminary heating of the precursor films to low temperatures below the 150 o C maximum working temperature of PET followed by 2) vigorous oxidation by treating the film with ozone in a UV Ozone (UV O 3 ) chamber In order to meet the low temperature processing requirements, we employed a precursor solution composed of nickel acetate tetrahydrat e and ethylenediamine ((NH 2 CH 2 CH 2 NH 2 en which created a solution composed mainly of [Ni(en) 2 (OAc)] + ions which decomp ose at lower temperatures, allowing oxidation of nickel. Combining this route to solution derived NiO at low temperatures with our low bandgap polymer poly dithienogermole thienopyrrolodione ( p DTG TPD), [6,23] we fabricated solar cells with an average PCE of 5.7 %. This performance is comparable to reference de vices with PEDOT:PSS as the HTL. The UV O 3 treatment of heated precursor films was examined by XPS, XRR and GIXRD. XPS analysis showed an increase in NiO and NiOOH upon exposure to ozone. XRR and GIXRD analysis showed an amorphous NiO film that is likely h ighly porous. T he solar cells with low temperature processed NiO were found to be slightly more air stable than devices with PEDOT:PSS under ambient conditions. Finally, as a proof of concept, solar cells with the low temperature processed NiO were fabrica ted on PET substrates, achieving an average PCE of 3.7%, promising for further development.

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86 3.2 Results and Discussion 3.2.1 Solar Cells from Low Temperature NiO Fabrication Part 1 In chapter 2, solar cells with NiO HTLs were fabricated by thermally decomposing films of a solution of nickel complexes, with the most prevalent ion in solution being [Ni(MEA) 2 (OAc)] + Films of this precursor solution partially decomposed at temperatures well below the 275 o C that formed optimized solar cells as shown in Figure 2 3 and Figure 2 10 Thus we hypothesized that the precursor transformation of the precursor into NiO could be completed at lower temperatures using other means than heat. In order to meet the maximum working temperature requirement of 150 o C of pol yethylene terephthalate (PET) substrates used in large scale, roll to roll (R2R) processing, we developed a two step transformation process to form NiO from the precursor film: 1) Partially decompose the precursor with heat and 2) finish the transformation into NiO using a reactive source of oxygen. For step 2, we employed ozone to since it is highly reactive and read ily made in a commercially available UV Ozone (UV O 3 ) chamber. The device structure for these sola r cells is the same as that of C hapter 2, an d is shown in Figure 3 1 As shown in Table 2 1, solar cells made from heating the NiO precursor solution to 185 o C showed little, but nonzero, photovoltage and photocurrent. Thus, we chose this working temperature as a starting point for the series solar cells with varying precursor decomposition temperatures. The solar cells were fabricated by heating the NiO precursor film to 185 o C, followed by exposing the films to ozone for an optimized time of 15 minutes. The same pDTG TPD/PC 70 BM BHJ active layer as used in C hapter 2 was spincast directly on the UV O 3 treated films We note that the UV O 3 treated pr ecursor films were darker than before the treatment, as discussed in more detail later The current voltage ( J V ) characteristics of the solar cell are sho wn in Figure 3 2 After UV O 3 treatment for 15 minutes, the solar cell drastically improved

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87 in all performance parameters. The V oc increased from 0.6 V to 0.82 V, the J sc increased from 0.09 mA/cm 2 to 12.8 mA/cm 2 the fill factor increased from 13% to 65%, and the PCE improved from 0.007% to 6.8%. These improved J V characteristics exhibited by this device are similar to devices made with PEDOT:PSS and the high temperature fabrication of NiO HTLs. The external qua ntum efficiency (EQE) shown in Figure 3 3 was also attractive The EQE remained above 50% between 435 nm and 695 nm. The integrated J sc from the EQE spectrum agrees to within a reasonable 11%. The difference is caused by spectral mismatch on the solar simu lator. While the precursor films fabricated with a precursor heating temperature of 185 o C and subsequent UV O 3 treatment generated high efficiency solar cells, proving the concept of a two step NiO formation, 185 o C is still too hot to be compatible with PET substrates in a R2R process. Thus we fabricated the same solar cells by using a first step temperature of 150 o C, the maximum working temperature of PET substrates. The J V characteristics are shown in Figure 3 4 Even after UV O 3 treatment of the precursor film, the solar cell photocurrent and photovoltage changed very little. The V oc remained nearly constant at 0.55 V, the J sc remained less than 0 .001 mA/cm 2 and the fill factor remained only 18%. Additionally, there was no visibl e film color change even upon prolonged UV O 3 treatment, indicating negligible chemical changes in the film. Therefore, we concluded that the chemical composition of this precursor solution was not appropriate for a low temperature fabrication of NiO. To r ealize low temperature fabrication of NiO compatible with R2R process temperatures, a new precursor solution was developed. 3.2.2 Low Temperature NiO Precursor In order to form NiO hole transport layers in solar cells at temperatures lower than the 150 oC limit of PET substrates, a new precursor solution was selected. In this solution, the coordinating amine, ethanolamine, was replaced by ethylenediamine (NH 2 CH 2 CH 2 NH 2 d ligand,

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88 it would donate electrons to nickel more strongly than MEA. Thus, we suspect that when the thermal decomposition of the precursor, which is actually th ermally assisted oxidation of Ni, took place, the ethylenediamine would be cleaved off the Ni a tom at lower temperatures The exact chemical oxidation route is unclear at this point and should be the focus of a future detailed study A chemical recipe that has been previously used to form NiO films [25] was modified for use here. This new precursor consists of nickel acetate tetrah ydrate (Ni(OAc) 4H 2 O) and ethylenediamine (NH 2 CH 2 CH 2 NH 2 ) in a 1:2 Ni 2+ :en mole ratio dissolved in ethanol. To determine the exact composition of the solution, we employed high resolution electrospray ionization mass spectrometry (ESI MS). As shown in Figure 3 5 The most common mononuclear ion in solution was found to be [Ni(en) 2 (OAc)] + with a mass to charge (m/z) ratio of 237.0786 and confirmed with theoretical modeling of the ion. Also shown in Figure 3 5 is the proposed chemical structure for this ion. For clarity, t he precursor solution based on the [Ni(MEA) 2 (OAc)] + ion will be referred to as Precursor A, while that containing [Ni(en) 2 (OAc)] + will be referred to as Precursor B. To ensure this new precursor would begin to decompose and form NiO at lower temperatures than 150 o C, films were spincast and heated to 140 o C. Additionally, these films underwent the same UV O 3 treatment as the films in the solar cells made from Precursor A. The optical absorption spectra are shown in Figure 3 6 The precur sor solution spectrum shows the typical d d transitions of octahedral nickel complexes with coordinating amines. Upon heating to 140 o C, many of the features of the precursor solution are weakened and broadened. The infrared absorption band is weakened and broadened, while the visible absorption band almost completely disappeared, and a band edge absorption in the UV appeared along with the remaining d d transition around 35 0 nm. Upon UV O 3 treatment of the heated film, much of the

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89 remaining precursor absor ption bands are eliminated, a broad visible absorption corresponding to Ni 3+ and the NiOOH surface species (details discussed later) appear, and a strong band edge absorption appears in the UV around 350 nm, corresponding to the NiO bandgap. 3.2.3 Solar Ce lls from Low Temperature NiO Fabrication Part 2 3.2.3.1 Device performance Since the precursor ions successfully decomposed and oxidized while only being heated to 140 o C, we fabricated solar cells in the manner previously described using these films with an optimized ozone treatment time of thirty minutes The J V characteristics of the solar cells are shown in Figure 3 7 The average V oc was 0.82 V, the J sc was 12.3 mA/cm 2 the fill factor was 57.5%, and the PCE reached 5.74%. The EQE spectrum is sho wn in Figure 3 8 The EQE was above 60% from 485 nm to 685 nm and reached a peak value of 64% at 610 nm. The integrated J sc obtained from the EQE spectrum agreed to within 5%. Thus, these devices are comparable in overall performance to the standard device s made from high temperature NiO processing and PEDOT:PSS, and are fabricated at temperatures amenable to using PET substrates in R2R fabrication. 3.2.3.2 Elucidating the UV O 3 treatment process To further investigate the role of the UV O 3 treatment in for ming NiO for these solar cells, we used X ray photoelectron spectroscopy (XPS) to track the chemical changes on the surface during the treatment. Films heated to 140 o C without any surface treatment were compared to those heated to 140 o C with surface trea tment for thirty minutes. Figure 3 9 shows the Ni 2p 3/2 spectrum before and after the UV O 3 treatment. The envelope can be deconvoluted into four individual peaks. For the films without any UV O 3 treatment, t he first peak centered at a binding energy of 854.2 eV corresponds to Ni O bonds in the standard octahedral bonding of bulk NiO [93 ,99,100] This is a reasonable assignment due to the

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90 partial decomposition and oxidation of the precursor as described above. The second large, peak centered at 855.6 eV is typically assigned to Ni OH bonds, especially those in Ni(OH) 2 or surface pass i vating hydroxyl bonds from airborne moisture exposure as well as Ni 3+ ions induced by nickel vacancies as discussed in Chapter 2 [93,97,99,100] However, it is more likely that a majority of the cont ribution to this peak comes from Ni coordinating with ethylenediamine and bonding to acetate groups remaining in the film as these low processing temperatures [98] The C 1s XPS spectrum shows several peaks, supporting the assertion of residual precursors in the film The third, very small peak, centered at 857.6 eV, has been associated wi th a surface dipolar oxyhydroxide (NiOOH) configuration [100] The NiOOH species has been shown to be a dipolar species capable of improving the performance of solar cells by increasing the effective work function of NiO when it was induced by oxygen plasma treatment [26,100] It is reasonable that all of these spe cies reside on the surface of the film simultaneously due to the heterogeneous nature of the partially decomposed and oxidized films. Finally, t he broad peak centered at 861.0 eV is assigned to shake up processes in the material structure [93] After UV O 3 treatment, the Ni O peak and NiOOH peaks grew substantia lly and are accompanied by a concomitant reduction in peak area for the precursor/ Ni OH peak. There is a corresponding intensity reduction in the C 1s spectrum after UV O 3 treatment indicating the leaving of carbon groups from the film due to the decompos ition The growth of the Ni O peak confirms that Ni is being oxidized by the ozone and forming NiO. The growth of the NiOOH peak confirms that this species is forming by ozone exposure as well Figure 3 10 shows the O 1s spectrum before and after UV O 3 tre atment. Before treatment, t he envelope can be deconvoluted into three peaks, corresponding to the three peaks in the main line of the Ni 2p 3/2 spectrum. The first peak, centered at 529.7 eV, is the companion to

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91 the Ni peak at 854.2 eV, and signifies Ni O o ctahedral bonding for standard bulk NiO [93,99,100] The large peak centered at 531 .4 eV is the companion to the Ni peak at 855.6 eV and signifies Ni OH bonding vacancy induced Ni 3+ [92,93,9 7,99,100] and can also have contributions from acetate bonding to Ni [98] The pe ak at 532.3 eV is the companion to the Ni peak at 857.6 eV, and signifies NiOOH [100] After UV O 3 treatment, a similar growth in the Ni O and NiOOH peaks occurs. This confirms the chemical changes shown in the Ni spectrum; namel y, growth of the Ni O structure, and increase in NiOOH centers. It should be noted that after the UV O 3 treatment, the films took on a dark black tint, signifying NiO OH and Ni 3+ based color centers [90,91,106,116] Similar surface chemistry changes have been seen while using oxygen plasma treatment on solution processed NiO [100] Additionally, a peak correspondin g to adsorbed water appears at 533.7 eV. This water may play a role in forming the surface NiOOH species that is essential for creating a favorable dipole in solar cells or may be a byproduct of the NiOOH formation Due to the low temperature growth of the films and the broad visible absorption rather than discrete visible absorption that would be caused by defect color centers, i t is likely th at the NiOOH dominat es Ni and O spectra rather than the vacancy induced Ni 3+ To confirm the extra contribution of water to the peak and confirm the lack of crystallinity that would allow a major contribution of the vacancy induced Ni 3+ structure to the XPS spectrum X ray reflectivity (XRR) and grazing incidence X ray diffraction (GIXRD) w ere used to monitor the formation of a layered NiO structure by the UV O 3 treatment. Figure 3 11 shows the XRR spectra of the films before and after th e surface treatment is applied. Before the treatment, the data show only one layer with a density of 2.23 g/cm 3 This is attributed to the plain heated precursor film. After the UV O 3 treatment, there were two layers present. The bottom layer had a density of 2.79 g/cm 3 while the top layer had a density of 1.08 g/cm 3 The

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92 bottom layer is then assigned to be the NiO formed by UV O 3 treatment, while the density of the top layer is nearly identical to water, confirming the water contribution to the high bindi ng energy peak in the O 1s XPS spectrum. The density of the NiO film is quite low compared to the bulk NiO density of 6.63 g/cm 3 This suggests that the NiO film structure formed is likely highly porous, defective, and amorphous. This is confirmed by exami ning the GIXRD spectrum, shown in Figure 3 12 The spectrum shows no discernable diffraction peaks, and only broad amorphous humps due to nearest neighbor interactions where the bulk diffraction peaks are calculated to exist. 3.2.3.3 Comparing low tempera ture NiO solar cells to standards In order to compare the performance of these solar cells to the standard cells with PEDOT:PSS HTLs and those using high temperature processed NiO. The J V characteristics of these three solar cells are compared in Figure 3 13 and Table 3 1 The V oc of the low temperature processed NiO based devices from Precursor B show comparable performance in all parameters to both the devices with PEDOT:PSS and high temperature processed NiO. The EQE spectra are also compared in Figure 3 14 The breadth and magnitude of the EQE for the low temperature processed solar cells are comparable to those of the high temperature NiO and the PEDOT:PSS devices. It is noteworthy that even though the NiO layer formed by UV O 3 treatment is completely amorphous and likely has residual precursor components in the film, the solar cells perform as well as the reference device with PEDOT:PSS, a well established hole transport material. This can be understood in terms of the strong surface dipole generated by the NiOOH species, which has been seen in other solar cells with NiOOH induced by oxygen plasma treatment of high temperature processed NiO [100] The surface dipole creates a large built in potential, generating a large inter nal electric field to extract charge carriers, even with an imperfect interface between the NiO and BHJ. Electroabsorption measurements shown in Figure

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93 3 15 show the built in potential to be 1.2 eV, more than 0.2 eV larger than that of the PEDOT:PSS device s and the untreated high temperature NiO devices. This excess built in field is essential to maintaining a high fill factor and short circuit current, and also plays a role in creating a large V oc in the devices 3.2.3.4 Air stability of the solar cells The final performance parameter important for determining if these novel solar cells are viable alternatives to the status quo is the air stability. Figure 3 16 compares the performance parameters of the three types of solar cells. The device with the new low temperature route to NiO is slightly more air stable than the device with PEDOT:PSS but less stable than the device with NiO fabricated from the high temperature process described in chapter 2 The degradation in performance for the low temperature Ni O based device occurs in all performance parameters, while it is mostly in the J sc for the device with PEDOT:PSS. Clearly the device and material degradation mechanisms are different. PEDOT:PSS has already been shown to be hygroscopic and its acidity is kn own to etch the ITO contact [65] In the device with PEDOT:PSS, the J sc loss is indicative of the contact degradation. The NiO was shown to contain a thin layer of water on the surface after air exposure. It is possible the s urface chemistry of the UV O 3 treated NiO changes over time, leading to various interfacial reactions as well as bulk chemical reactions which adversely affect the device performance. 3.2.4 Solar Cells Fabricated on Plastic Substrates Since the ultimate, long term goal, of this study was to develop a method for low temperature processing of NiO interfacial layers in solar cells that are compatible with PET substrates, we tested the robustness of the fabrication process by fabricating full solar cells on PET. ITO anodes were sputtered onto clean PET sheets, and the entire fabrication process was conducted th e same as on glass substrates. A photograph of the device is shown in Figure 3 17

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94 and the J V characteristics are shown in Figure 3 18 The device on PET shows an average V oc of 0. 8 0 .01 V, J sc of 9.9 0.3 mA/cm 2 fill factor of 4 6.5 0.5%, and PCE of 3.7 0.1 %. This PCE is comparable to the community standard small area devices made on glass in well established fabrication methods with a P3HT:PCBM BHJ active layer with PEDOT:PSS hole transport layer s Thus, the a chievement of nearly 4% PCE is promisi ng for future applications of these materials, and this fabrication method for NiO in particular. 3.3 Conclusions In summary, we have discovered a novel route to fabricate NiO hole transport layers in polymer photovoltaic cells at working temperatures low enough to be compatible with PET substrates. A two step procedure for fabricating NiO at these low temperatures was developed. The precursor film was partially decomposed with heat, followed by an extended oxidation period under ozone exposure in a UV O 3 chamber. During the ozone treatment phase, NiO and NiOOH are formed, as monitored by XPS. The NiOOH acts as an interface dipole which aids in hole extraction and electron blocking. NiO films fabricated from this technique are completely amorphous and high ly porous. Solar cells incorporating this NiO as an anode interfacial layer, or hole transport layer, benefit from the NiOOH dipole, generating a PCE of 5.7%, comparable with a reference device fabricated with PEDOT:PSS as the HTL. The devices containing t his novel NiO layer are slightly more stable in air than those with PEDOT:PSS. Finally, as a proof of concept, solar cells with this low temperature NiO were fabricated on PET substrates which achieved an average PCE of 3. 7%, promising for future developme nt. 3.4 Experimental Procedure NiO precursor solutions. Precursor A: Nickel acetate tetrahydrate (Ni(CH 3 COO) 2 4H 2 O) (Acros Organics) was dissolved in ethanol with monoethanolamine (NH 2 CH 2 CH 2 OH) (Sigma Aldrich) (0.1 mol L 1 ). The mole ratio of Ni 2+ : MEA was maintained at 1:1 in solution.

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95 Dissolution took place while stirring in a sealed glass vial under air at 70 C for four hours. The solution appeared homogeneous and deep green after approximately forty minutes. Precursor B: Nickel acetate tetrahydrate (Ni(CH 3 COO) 2 4H 2 O) (Acros Organics) was dissolved in ethanol with ethylenediamine (NH 2 CH 2 CH 2 NH 2 (en) ) (Sigma Aldrich) (0.1 mol L 1 ). The mole ratio of Ni 2+ :en was maintained at 1:2 in solution. Dissolution took place while stirring in a sealed glass via l under air at 70 C for four hours. The solution app eared homogeneous and purple after approximately forty minutes. NiO film preparation All NiO films were deposited by spin casting onto the appropriate substrates. For solar cells, the substrates were IT O coated glass with sheet resistance of 15 ohm sq 1 For GIXRD, the substrates were plain glass. For XPS, the substrates were single crystal Si wafers. All substrates were cleaned by the following procedure before spincasting: successive sonication in DI w ater, acetone, and isopropanol baths for 15 minutes each. NiO films were deposited by spincoating the precursor solution (B) onto cleaved Si wafers, followed by heating to 140 o C for 45 minutes. Samples that underwent UV O 3 treatment were treated for 30 minutes in a UV Ozone chamber. The UV Ozone cleaner is a Jelight UVO Cleaner Model 42 with wavelengths of 254 nm and 185 nm with intensity of 30 mW/cm 2 Device f abrication and c haracterization To fabricate the solar cells, I TO coated glass substr ates were UV ozone treated for 2 5 minutes between solvent cleaning and spin casting of the Precursor B solution. No UV Ozone treatment was necessary to wetting of Precursor A or PEDOT:PSS. PEDOT:PSS (Clevios Al 4083 ) was filtered thro ugh a 0.45 m PTFE filter prior to spin casting at 7000rpm. The PEDOT:PSS films were heated at 140 o C for 10 minutes in air and transferred immediately to the nitrogen filled glovebox for further processing. NiO films were deposited by spincoating the prec ursor solution (B) onto substrates, followed by heating to 140

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96 o C for 45 minutes. Subsequently, UV O 3 treatment was carried out for 30 minutes in a UV Ozone chamber. When NiO was used as the HTL instead of PEDOT:PSS, the NiO films were allowed to cool befo re being transferred to a nitrogen filled glovebox for further processing. The donor polymer used was poly dithienogermole thienopyrrolodione ( p DTG TPD) and the acceptor fullerene was [6,6] phenyl C 71 butyric acid methyl ester (PC 71 BM) The polymer:fullere ne blend ratio is 1:1.5 by weight and is dissolved in chlorobenzene with 5 vol% 1,8 diiodooctane (DIO) additive. PC 71 BM was acquired from Solenne (purity >99%) and was used as received. The blend was spin cast on top of the HTL film to give an active layer thickness of 105 nm. The devices were immediately loaded into a thermal evaporator where 1 nm of LiF and 100 nm of Al were deposited under a vacuum pressure of 5x10 7 Torr. For devices using PET substrates, ITO was sputtered in a Lesker PVD 75 DC sputtere r to a thickness of 90 nm using an anode shadow mask. The device area was 0.046 cm 2 Current voltage ( J V ) characterization was performed with a Keithley 4200 semiconductor parameter analyzer system with a Newport Thermal Oriel 94021 1000 W solar simulator ( 4 in. by 4 in. beam size) using the AM1.5 G solar spectrum at 100 mW cm 2 incident power The light intensity was calibrated by a n ORIEL 91150V monosilicon reference cell calibrated by Newport Corporation. EQE measurement s w ere conducted using a n in hous e setup consisting of a Xenon DC arc lamp, a n ORIEL 74125 monochromator, a Keithley 428 current amplifier, a n SR 540 chopper system and a n SR 830 DSP lock in amplifier from SRS The experimental setup for electroabsorption was described previously. [112] GIXRD and XRR measurements GIXRD measurements were performed on a Philips XPert MRD with Cu K X rays (non monochromatic) with incident angle = 3. XRR was also carried out on the Philips XPert MRD with Cu K X rays (non monochromatic).

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97 XPS measurements XPS data were acquired on a Perkin Elmer 5100 XPS system with a non monochromatic Al anode X ray source with a pass energy of 35.75 eV Peak d econvolution and analysis was performed in RBD AugerScan version 3.2 software. The adventitious C 1s peak was referenced to 284.8 eV. Mass s pectrometry measurements Solutions were introduced into an Agilent 6210 TOF MS via direct injection followed by ele ctrospray ionization (ESI) with an autosampler. The mobile phase was spectroscopy grade ethanol from Fisher Scientific and was used as received. Accurate mass identification was pe rformed in MassHunter software. 3.5 Figures and Tables Figure 3 1. Device structure and energy levels with respect to vacuum of the materials used in the photovoltaic cells.

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98 Figure 3 2. Current voltage characteristics of solar cells fabricated with Precursor A for NiO, heated to 185 o C, then UV Ozone treated to complete the formation of NiO. A drastic increase in all device performance parameters is observed after UV Ozone treatment of the heated films. Figure 3 3. External quantum efficiency spectrum of solar cells fabricated with Precursor A for NiO, heated to 185 o C, the n UV Ozone treated to complete the formation of NiO.

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99 Figure 3 4. Current voltage characteristics of solar cells fabricated with Precursor A for NiO, heated to 150 o C, then UV Ozone treated to complete the formation of NiO. At such a low temperature, the precursor was insufficiently decomposed and Ni was hardly oxidized. The resulting solar cells show negligible photocurrent and photovoltage output. Figure 3 5. High resolution mass spectrum showing the isotopic distribution for the [Ni(en) 2 (OAc)] + ion. The main peak for Ni 58 is located at m/z = 237.0786.

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1 00 Figure 3 6. Absorption spectrum showing the progression of the NiO film fabrication by the low temperature route for Precursor B. Figure 3 7. Current voltage characteristics of the solar cells with NiO fabricated from Precursor B with the low temperature process.

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101 Figure 3 8. External quantum efficiency spectrum of the solar cells with NiO fabricated from Precursor B with the low temperature process.

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102 Figure 3 9. High resolution deconvolution of the Ni 2p 3/2 XPS spectrum before and after UV Ozone treatment. The envelope shows four peaks, one for NiO (dark red) one for Ni OH bonding precursor coordination or defect induced Ni 3+ (blue), one for NiOOH (light green) and a broad shake up peak at the high binding energy (blueish green) After UV Ozone treatment, the NiO and NiOOH peak increase, signifying the formation of NiO and NiOOH, with a concurrent decrease in Ni OH/ precursor coordination peak.

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103 Figure 3 10. High resolution deconvolution of the O 1s XPS spectrum before and after UV Ozone treatment. The envelope shows three peaks, one for NiO (dark red) one for Ni OH bonding/ precursor coordination / defect induced Ni 3+ (blue), and one for NiOOH (blueish green at 532.3) After UV Ozone treatment, t he NiO and NiOOH peak increase, signifying t he formation of NiO and NiOOH with a concurrent decrease in Ni OH/precursor coordination peak. Additionally, a water peak appears after UV Ozone treatment (blue green at 534 eV) After UV Ozone treatment, the Ni OOH peak is shown in bright green.

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104 Figure 3 11. X ray reflectivity (XRR) spectrum of the precursor films before and after UV Ozone treatment. Curve fitting reveals that before UV Ozone treatment, there exists only one layer, while after UV Ozone treatme nt, there exists two layers one is the highly porous and defective NiO with a top layer of water, confirming the contribution in the XPS spectrum.

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105 Figure 3 12. Grazing incidence X ray diffraction spectrum showing a completely amorphous NiO film after the UV Ozone treatment. Figure 3 13. Current voltage characteristics of the solar cells fabricated with the low temperature NiO process compared to those of solar cells fabricated with PEDOT:PSS or high temperature NiO.

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106 Figure 3 14. External quantum ef ficiency spectrum of the solar cells fabricated with the low temperature NiO process compared to those of solar cells fabricated with PEDOT:PSS or high temperature NiO. Figure 3 15. Electroabsorption measurement showing a built in potential of 1.2 eV in the solar cells with the low temperature processed NiO HTL. The large built in potential is due to the NiOOH surface dipole created during the UV Ozone treatment formation of the NiO film.

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107 Figure 3 16. Air stability of the solar cells with low temperature processed NiO compared to devices with PEDOT:PSS and high temperature processed NiO as references. The devices with low temperature processed NiO are slightly more air stable than those with PEDOT:PSS but less stable than those with high temperature processed NiO.

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108 Figure 3 17. Photograph of the solar cells fabricated with low temperature NiO on PET substrates. Photo co urtesy of Jesse Mande rs Figure 3 18. Current voltage characteristics of the solar cells fabricated on PET substrates. The devices achieved an average PCE of 3.7%, promising for further development.

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109 Table 3 1. Performance of solar cells fabricated with PEDOT:PSS, high temperature processed NiO, and the low temperature processe d NiO from Precursor B. One standard deviation is reported in parenthesis. HTL V oc (V) J sc (mA cm 2 ) FF (%) PCE (%) V bi (EA) PEDOT:PSS 0.83 (0.005) 11.9 (0.5 ) 58.5 (0.7 ) 5.8 (0.1) 0.96 V NiO 275 o C 0.82 (0.006) 13.9 (0.3) 68.4 (0.4) 7.8 (0.2) 0.95 V NiO Precursor B, 140 o C + UVO 3 0.82 (0.01) 12.3 (0.3) 57.5 (0.5) 5.7 (0.1) 1.2 V

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110 CHAPTER 4 AIR STABLE MULTISPECTRAL PHOTODETECTORS WITH LOW NOISE MADE FROM ALL SOLUTION PROCESSED INORGANIC SEMICONDUCTORS 4 .1 Background Infrared (IR), visible (Vis), and multispectral photodetection systems are widely used in optical communications, imaging, security, ranging, and household electronics. Due to the prevalence of these systems in modern society, high quality photodetectors made from inexpensive an d high throughput manufacturing techniques are in high demand. Currently, most photodetectors for these applications are photodetectors for these applications are grown by epitaxial deposition techniques, then bonded to read out integrated circuits (ROICs) in a lattice of individual sensor pixels known as a focal plane array (FPA). This process can be costly and usually only allows for small area device fabrication. To mitigate these issues, solution processed semiconductors serve a key role in emerging pho todetector technologies as they can be processed quickly on large area substrates by roll to roll processing and do not necessarily need to meet stringent vacuum or lattice matching constraints. Solution processed devices made from colloidal quantum dots (QDs) are promising for use in the near infrared (NIR), short wave infrared (SWIR), mid wave infrared (MWIR) and visible regions [28,44] QDs are advantageous because their absorption spectrum can be tuned easily with the quantum confinement effect controlling the size of the QDs during synthesis [33,34] Specifically, lead sulfide (PbS) QDs hav e been used as the photoactive layer in many photodetectors [57 61,117] due to their spectral tunability across the NIR and visible spectrum relative ease of synthesis and relatively large exciton Bohr radius of 18nm [52] The Sargent group first reported PbS photodetectors with a PbS/MEH PPV polymer blend photoactive layer in a photodiode structure [118] Later, elegant ligand exchange procedures allowed lateral channel photoconductors to be developed that achieved extremely high

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111 responsivity, and detectivity comparable to commercially available photodetectors [58] While photoconductors offer strong ph otoresponse, they are typically much slower than photodiodes due to their trap induced gain mechanisms [28,117] To solve this issue, Schottky junction photodiodes were developed with MHz bandwidth but the devices achieved lower D* compared to the slower photodiodes and photoconductors that had higher photoresponse [59] More recently, the Klimov group fabricated PbS/ZnO and PbS/TiO 2 p n junction photodiodes that exhibited low noise levels, and favorable D* and LDR. However, the best performing device with ZnO as the n type layer incorporated the organic hole ex traction layer PEDOT:PSS, which is known to be unstable in air [65,115] There has yet to be a PbS QD photodetector reported that simultaneously achieves D* > 1x10 12 large bandwidth in the kHz, large LDR, low noise, and long air lifetime. In a vertically stacked P I N like photodiode, charge blocking layers are inserted between the electrodes and the photoactive layer to reduce the dark current ( J d ) and thus the noise in the device. Previously, colloidal ZnO nanocrystals were used as a wide bandgap hole blocking layer/electron transport layer (HBL/ETL) with an organic electron blocking layer (EBL/HTL) in PbSe based double heterojunction devices [62] This combination significantly lowered dark current and improved the air stability of the device. However, there was still room for improvemen t to lower the dark current and hence the noise, while improving the air stability and photosensitivity of the device. In addition to colloidal ZnO, which has been shown to serve as an air stable, n type HBL/ETL, we chose NiO to serve as the p type oxide HTL/EBL. We have previously shown that solution derived NiO is an air stable wide bandgap p type semiconductor which can transport holes and can be used as an electron blocker [115] The use of both of these wide bandgap oxides as carr ier blocking layers creates a novel all inorganic double heterojunction which reduces the

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112 dark current, while encapsulating the PbS QDs layer from moisture and harmful UV irradiation. Therefore, we expect that the PbS QDs devices with these blocking layers should have good air stability and a broad spectral responsivity with low noise. Here, we show that the double heterojunction structure does indeed enable low noise equivalent power (NEP) of tens of picowatts across the working spectrum and a high specif ic detectivity ( D* ) value on the order of 1x10 12 cm Hz 1/2 /W (1 cm Hz 1/2 /W = 1 Jones). The noise levels are half an order of magnitude lower than the lowest reported [61] multispectral photodio de noise levels, more than an order of magnitude lower than the lowest multispectral photoconductor noise levels [58] The D* values are similar to those of commercially available photodetectors made from Ge, InGaAs and Si. Additionally, due to the presence of the oxide layers, even without encapsulation, the resulting devices have extremely good air stability with storage lifetime up to five months with no degradation in device performance. This is the longest lifetime reported for any photodetector with a pure PbS QD photoactiv e layer. 4.2 Materials Properties 4.2.1 PbS Quantum Dots PbS QDs were synthesized using the hot colloidal injection technique which has been shown to yield highly air stable QDs [119] Figure 4 1 shows a typical absorption spectrum of PbS QDs synthesized with this route. Along with the as synthesized spectrum, the absorption spectrum of the same batch stored in air for three weeks is included. The spectrum is unchanged, showing the robustness of the synthesis and air stability of the quantum dots. Figure 4 2 shows the transmission electron micrographs of the QDs showing quasi spherical c rystals of approximately 4 5 nm in diameter. During device fabrication 4.5 nm diameter quantum dots were used with an absorption edge of approximately 1.3 m. During the deposition of the QD layer, a solid state ligand exchange was carried out to replace the native long insulating oleic

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113 acid ligands with shorter 1,3 benzenedithiol ligands. This exchange procedure decreases the spacing between the QDs in the film and has been shown to in crease the QD film conductivity, which improves the photoresponse of th e photodetectors [36] A schematic of the physical changes that take place during the ligand exchan ge in the QD film is shown in Figure 4 3. Complete fabrication details are provided in the Experimental Procedure section. 4.2.2 Zinc Oxide Nanoparticles Zinc oxide nanoparticles were synthesized by dropwise addition of a stoichiometric amount of a solution of tetramethylammonium hydroxide (TMAH) in ethanol to a solution of zinc acetate dihydrate dissolved in dimethyl sulfoxide (DMSO) under continuous stirring. After precipitation and washing, the nanoparticles were dispersed in pure ethanol. The zinc oxide nanoparticles are nearly spherical with a diameter of approximately 6 nm as shown in the transmission electron micrograph in Figure 4 4 The 6 nm nanoparticle diameter was confirmed by applying the Scherrer equation [102] to the X ray diffraction spectrum shown in Figure 4 5 The diffraction spectrum matches that of ZnO in the typical hexagonal wurtzite crystal structure. The optical absorption spectrum of the ZnO nanoparticles is shown in Figure 4 6 The nanoparticles have an abs orption edge of 36 5 nm, c orresponding to a bandgap of 3.4 eV making them transparent to the visible and IR radiation that the PbS absorbs. 4.2.3 Nickel Oxide Films Nickel oxide films were prepared using the same method as described in C hapter 2. The device optimization procedure for these devices led to 500 o C being the optimum NiO formation temperature creating a polycrystalline NiO film, as shown in the grazing incidence X ray diffraction spectrum shown in Figure 4 7 The transmission spectrum of the NiO film fabricated for these devices is shown in Figure 4 8 The spectrum shows transmission greater tha n 95% across the entire visible and NIR spectra

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114 4.3 Photodetector Performance In Figure 4 9 we show the structure of our photodetectors. The PbS QD layer is sandwiched between a ZnO ETL/HBL and a NiO HTL/EBL to form a double heterojunction. Also, since ZnO and NiO are wide bandgap semiconductors, as shown in the energy band diagram in Figure 4 10 they are transparent to visible and infrared radiation and ideal for carrier transporting and blocking layers. 4.3.1 External Quantum Efficiency and Responsivity The two basic parameters of photodetector sensitivity are the external quantum efficiency (EQE) and the responsivity. Figure 4 11 shows the responsivity and EQE at 1 V applied bias. The EQE in these devices reached 24% at 1130 nm and 52% at 575 nm while the responsivity reached 0.2 A/W at 1130 nm and 0.25 A/W at 600 nm. These responsivity valu es are on the same order of magnitude as the values reported for commercially available photodetectors. The broad spectral response of these photodetectors makes them useful for multispectral applications. 4.3.2 Speed and Bandwidth To evaluate these photo detectors for imaging applications, we determined their temporal response. Typical video frame rates are 30 Hz, so a photodetector used for these purposes would need a bandwidth greater than 30 Hz to ensure good image quality. To measure the speed of our d evices we used a pulsed LED light source and measured the rise and fall times of the photocurrent on an oscilloscope. Details of the experimental setup are described in the Methods section. The temporal response data for light intensity of 54 mW/cm 2 at 410 nm for different bias voltages are shown in Figure 4 12 The temporal response of the device was symmetrical, with characteristic time constants on the order of a few microseconds. Using the temporal response, we calculated the bandwidth from the relation

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115 (4 1) where B is the bandwidth in Hz, and is the characteristic time constant of the device. F igure 4 13 shows the bandwidth as a function of applied bias. The bandwidth of the device increased linearly with applied bias. Furthermore, the bandwidth is dependent on the incident light intensity, which indicates that the speed of the photodetector is transport limited. The light intensity dependence of the bandwidth is shown in Figure 4 14 This light intensity dependence has been shown in ot her QD based photodetectors and can be explained with a trap filling model [57,59] as follows. Under low intensity excitation, photo generated carriers are affected by the large trap density of the QD film, resulting in a slow response. At higher intensity excitation, a significant portion of the traps are filled by photo excited carriers, and the carrier transport becomes trap free, allowing a faster photoresponse. This model is further confirmed in these photodetectors by the sharp increase in bandwidth at high light intensities. The bandwidth at high light intensity reached 36 kHz. Furthermore, even at low light intensities comparable to bright night skies, the detector bandwidth was 4 kHz, orders of magnitude faster th an the requirements for conventional video frame rates. 4.3.3 Noise Levels and Detectivity since it is related to the lowest level signal that can be captured by the device. There are several types of noise in The power spectral density of pink noise falls as 1/f, while that of white noise is constant in frequency space.

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116 The causes of 1/f noise in semiconductors are still debated and no closed form solution unifying all proposed causes exists. However, among t he most common explanations are time varying conductivity via mobility fluctuations and carrier density fluctuations [120,121] Trapping and detrapping of charges act as generation recombination processes that change the carrier density [122,123] while scattering events along surfaces and inter faces [120] or in the bulk [124,125] act to change the carrier mobility. In re ality, several of these mechanisms may be acting simultaneously [126] There is no known experimental lower frequency limit to 1/f noise, but there exists an upper limit at which the noise falls below the other type of noise present shot noise. The transition frequency between the 1/f noise dominated frequency space and the shot noise dominated frequency space is called the corner frequency ( f c ). In general, the corner frequency is a signifier of the quality of transport and physical structure of the device; since 1/f noise is higher than shot noise at f < f c it is advantageous to hav e the lowest possible f c to realize a quiet device. At frequencies above the corner frequency, shot noise dominates. Shot noise is the random fluctuation in the DC dark current due to the quantization of electrons and represents the noise floor of this cla ss of photodetectors. Specifically, the inherent quantization of electrons causes them to arrive at the external circuit with randomly distributed time intervals between ent fluctuation ( i.e noise) that is independent of frequency. Because shot noise is independent of frequency, there is no upper frequency limit. So at f > f c shot noise dominates through the entire bandwidth of the device. This is important in the D* cal culation as discussed later. The noise spectrum was measured in a series of reverse biases to obtain a more complete understanding of the device, as shown in Figure 4 15 At high biases, 1/f noise was prominent at

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117 low frequencies. As the bias decreased, DC dark current ( J d ) decreased and 1/f noise decreased more than the shot noise. This is because the 1/f noise can go as the J d 2 in heterojunction devices [127] while shot noise is usually linearly proportional to J d As the bias decreased and 1/f noise decreased, f c decreased from ~70 Hz at 3 V and ~30 Hz at 2.5 V to less than 1 Hz at 1 V. At biases of 1 V and below, the entire measured noise spectrum is dominated by shot noise. The cutoff for 0.5 V at 10 Hz is caused by the narrow bandwidth on the extremely high gain amplifier setting needed for the low dark current. These corner fre quencies are among the lowest reported for PbS QD photodetectors in the literature and are indicative of clean transport through the device since 1/f noise is minimized. This is an advantage of device structures with vertical transport. Since conduction in the device is vertically through the bulk, the unwanted noise contribution of carrier trapping and scattering at surfaces and interfaces that can appear in the 1/f regime [120,123] for lateral devices is minimized. The low overall noise levels are enabled by low J d in the device approximately 20 nA/cm 2 at 0.5 V and 34 nA/cm 2 at 1 V, without any device cooling, as shown in the current voltage plot in Figure 4 16 The device area was 0.046 cm 2 The f c and overall noise levels are among the lowest reported in the literature for PbS QD based devices and are a testament to the double oxide heterostructure. The absolute noise level provides the noise equivalent power (NEP), or the input power at which the s ignal to noise ratio is unity. The NEP is the last parameter of the D* calculation, The NEP is commonly calculated in the literature according to Eq uation 4 2 as follows [61,128,129] : (4 2)

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118 where < i 2 > is the integrated mean square noise spectral density in the device bandwidth and R is the sp ectral responsivity. Then, specific detectivity, D* can be calculated with Eq uation 4 3 below: (4 3) where A is the device area, B is the bandwidth, and NEP is the noise equivalent power. Using these relations, we found the NEP at 1V to be on the order of tens of pW across Figure 4 17 When the NEP is normalized to a 1 Hz bandwidth, it falls to tens of fW/Hz 1/2 on the order of reported values for commercial Si and InGaAs photodetectors. When D* is calculated, as shown in Figure 4b, values of 1.1x10 12 cm Hz 1/2 /W at 1135nm and 1.2x10 12 cm Hz 1/2 /W at 600nm are obtained. Also shown in Figure 4c are average values for the peak D* found in commercially available photodetectors. It is worthwhile to note that the D* for our PbS QD photodetector is within an order of magnitude of these detectors and already exceeds the D* found in many Ge photodetectors. Also, these other photodetecto rs are typically fabricated using expensive and lower throughput epitaxial deposition systems. In particular, high quality InGaAs photodetectors are typically made in MBE or MOCVD reactors, must be lattice matched to their substrate, and can have complicat ed bonding to the ROIC. Thus, InGaAs FPAs can cost tens of thousands of dollars. The only major limitation on a FPA with these PbS QD detectors is the ROIC geometry, since the active materials could be deposited from solution directly onto the ROIC. While the D* of our PbS photodetectors is similar to that of commercial photodetectors, another parameter is needed for imaging bright and dark objects in the same frame without losing image quality: the linear dynamic range (LDR ).

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119 4.3.4 Linear Dynamic Range Ano ther device parameter important to a photodetector is the linear dynamic range (LDR) which can be calcula ted using Equation 4 4 : (4 4) where P max is the highest incident light power and P min is the lowest incident lig ht power of the range in which the photodetector response is linear with the incident power. Figure 4 18 shows the measured LDR of a typical PbS photodetector. The measured LDR extended 67 dB, comparable to the LDR of InGaAs photodetectors. [56] 4.3.5 Air Stability Air stabil ity data for QD based photodetector devices are not often reported in the literature. However, previous studies with these oxide materials used in this device are promising in terms of device lifetime [62,115] It should be noted that the PbS layer here is only protected by the two oxide layers in this double heterostructure and the device stability is expected to be good. Figure 4 19 shows the device air stability data over a period of 5 months for a device without air at room temperature. In the first few days of storage, the device perfor mance actually improved slightly and there is no degradation over the entire testing period. This initial increase has been seen in other devices with ZnO [62,130,131] Since colloidal ZnO has a large surface defect density [132] atmosp heric or residual oxygen and moisture can easily bind to the surface of the nanocrystals and not penetrate through the film into the other layers. Also, photoexcitation [16] and passivation of the surface defects [133] can enhance the conductivity of ZnO, leading to an improvement in the device. Another advantage of ZnO is its protection from cathode oxidation. Since aluminum cathodes are prone to oxidation, dir ect contact between the cathode and the photoactive layer or oxidation prone organic blocking layers causes stability problems [62]

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120 However, when aluminum is depos i ted on ZnO, as in the photod etectors in this report, any interfacial mixing or oxidation that takes place would not hinder the device performance because aluminum zinc oxide is an air stable n type transparent conductor [134] In addition, NiO has also been shown to improve the device stabi lity when it replaced organics that are susceptible to moisture and oxygen attack. The presence of the NiO layer here also offers additional protection to the device. The air stability lifetime shown here is the longest reported in the literature for any P bS QD based photodetector. 4.4 Conclusions In summary, we have shown our high quality PbS QDs with effective ligand passivation PbS QDs were incorporated into a photodetector with solution processed, air stable oxide charge blocking layers. The resulting device has a wide spectral photoresponse fro m the NIR region through the entire visible spectrum. The bandwidth of these devices is sufficient for high frame rate imaging applications at low light intensities. The NEP, D* and LDR values are comparable to commercially available photodetectors made f rom expensive vacuum based methods. Finally, the solution processed air stable oxides act as carrier transport and blocking layers, and the initial air stability study indicates these device are extremely robust with long lifetime. 4.5 Experimental Procedu re PbS s ynthesis 0.444g of PbO (Sigma Aldrich) was dissolved with 2 mmol of oleic acid in a three necked flask under Ar flow with 55 mL of octadecene (ODE). The solution was heated and held at 120 o C. 210 L of hexmethyldisathiane ((TMS) 2 S) (Sigma Aldrich ) was dissolved in 5 mL of ODE and injected into three necked flask. The reaction continued for four minutes. The reaction was quenched and precipitated with cold acetone. The QDs were washed and redispersed three times with acetone and chloroform. The fin al product was stored dry until use.

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121 TEM images of PbS were taken on a JEOL 2010F with an accelerating voltage of 200 kV at the University of Florida Major Analytical Instrumentation Center (MAIC). ZnO s ynthesis The synthesis has been previously published [16,133] The synthesis was performed by dropwise addition of a stoichiometric amount of tetramethylammonium hydroxide (TMAH) (0.55 M) to 30 mL of 0.1 M zinc acetate dihydrate dissolved in dimethyl sulfoxide (DMSO) under continuous stirring. After precipitation and washing, the nanoparticles were dispersed in pure ethanol. TEM images of ZnO were taken on a JEOL 2010F with an accelerating voltag e of 200 kV at the University of Florida Major Analytical Instrumentation Center (MAIC). NiO precursors The procedure was previously published for use in solar cells [115] A 0.1 M solution of nickel acetate tetrahydrate in ethanol was made. A 1:1: mole ratio of monoethanolamine (MEA) to nickel was added as a complexing agent. Solution was stirred until all reagents dissolved into a green solution. Device f abrication The NiO precursor solution was spincast onto solvent and UVO 3 cleane d ITO coated glass substrates and heated to 500 o C in air for one hour to form continuous NiO films. After cooling, the substrates were transferred to a nitrogen glovebox for the PbS layers. PbS layers were spincast from a dilute suspension of QDs in chlor oform. After each layer deposition, the films were soaked in a 1.0 M solution of 1,3 benzenedithiol in acetonitrile for the ligand exchange. This PbS film deposition and ligand exchange was repeated to yield approximately 250nm thick films. ZnO nanoparticl es were directly spincast on top and the device was heated to 80 o C for 15 minutes. Then a 100 nm thick aluminum cathode was thermally evaporated at chamber pressures of ~10 6 Torr. Device c haracterization All characterization and noise measurements were p erformed at room temperature. Current voltage ( J V ) characterization was performed with a Keithley 4200

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122 semiconductor parameter analyzer system. EQE and Responsivity measurement s w ere conducted using a n in house setup consisting of a Xenon DC arc lamp, a n ORIEL 74125 monochromator, a Keithley 428 current amplifier, a n SR 540 chopper system and a n SR 830 DSP lock in amplifier from SRS Bandwidth o ptical measurements were conducted with an in house pulsed LED setup (Thorlabs) The setup consists of the LED driven by a function generator with a square voltage high bandwidth mode and read by a digital oscilloscope on one channel, while t he LED driver was monitored by another channel. Linearity measurements were confirmed by using neutral density filters measured by a Perkin Elmer UV Vis spectrophotometer, coupled with broadband solar spectrum light from Newport Thermal Oriel 94021 1000 W solar simulator ( 4 in. by 4 in. beam size) using the AM1.5 G solar spectrum at 100 mW cm 2 incident power. The light intensity was also calibrated by an ORIEL 91150V monosilicon reference cell calibrated by Newport Corporation. Noise m easurement : Devices w ere packaged and bonded inside a small aluminum Faraday cage and placed inside a solid copper Faraday cage with a SR570 current preamplifier (electrically isolated) which acted as the bias source and amplifier. The solid aluminum and copper boxes shielded the device and amplifier from outside noise. The SR570 amplifier was then connected to an Agilent 35670A Dynamic Signal Analyzer spectrum analyzer. The spectrum analyzer output the signal directly to the data acquisition setup. A modified setup and equival ent circuit was published previously [128]

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123 4.6 F igures Figure 4 1. Typical absorption spectrum of PbS QDs synthesized with the hot colloidal injection technique. The same batch stored in air was measured three weeks later and showed an unchanged absorption spectrum. This indicates the QDs are well cap ped by the oleic acid ligands. Figure 4 2. Transmission electron micrographs of PbS QDs used in the photodetectors. The scale bar is 20 nm on the main image and 10 nm on the inset The image shows nearly spherical quantum dots with appr oximately 2 nm in terdot spacing, corresponding to the oleic acid ligand length.

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124 Figure 4 3. Schematic of the change in spacing between quantum dots upon exchanging oleic acid for 1,3 Benzenedithiol. The shorter benzenedithiol ligands decrease the spacing between quantum dots. Figure 4 4. Transmission electron micrographs of 6 nm ZnO nanoparticles used in the PbS photodetectors. The scale bar is 50 nm on the main image and 5 nm on the inset.

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125 Figure 4 5. X ray diffraction spectrum of the ZnO nanoparticles used in the PbS photodetectors. The diffraction spectrum confirms that the nanoparticles are 6 nm in diameter and have the Wurtzite crystal structure. Figure 4 6. Optical absorption spectrum of the ZnO nanoparticles used in the PbS photodetectors. The absorption edg e is 365 nm, corresponding to a bandgap of 3.4 eV.

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126 Figure 4 7. NiO diffraction spectrum showing polycrystalline NiO in the typical rocksalt crystal structure. Figure 4 8. Transmission spectrum of NiO films used in the PbS QD photodetector. The transmis sion spectrum shows that the films are greater than 95% transparent through the entire visible and NIR spectrum.

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127 Figure 4 9. The physical device structure for the photodetectors. Figure 4 10. Energy band diagram for the photodetector. The energy band alignments create a favorable double heterojunction structure.

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128 Figure 4 11. Plot of responsivity and EQE for the photodetectors. EQE in these devices reached 24% at 1130nm and 52% at 575nm while the responsivity reached 0.2 A/W at 1130nm and 0.25 A/W at 600nm. Figure 4 12. Temporal response of the photodetector for light intensity of 54 mW/cm2 at 410 nm.

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129 Figure 4 13. Speed and bandwidth as a function of applied reverse bias. The bandwidth increases with applied bias. Figure 4 14. Photodetector bandwidth as a function of incident light intensity at 1 V applied bias. As light intensity increases, photoexcited carrier density becomes sufficient to fill traps, allowing an increase in bandwidth.

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130 Figure 4 15. Noise current spectral density as a function of frequency. At high applied biases, 1/f noise appeared with low f c values. As bias is decreased, J d decreased and 1/f noise is pushed to lower frequencies than measured in our experiment. At 1 V applied bias, the whole bandwidth is shot noise dominated. Figure 4 16. Dark current density vs. applied bias. This plot shows extremely low dark current enabled by the novel double heterostructure.

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131 Figure 4 17. Noise Equivalent Power and Specific Detectivity of the photodetectors at 1 V. NEP D* was calculated to be as high as 1.1x10 12 cm Hz 1/2 /W at 1135nm and 1.2x10 12 cm Hz 1/2 /W at 600nm. For reference, reported maximum values for D* in commercial photodetectors are indicated. Figure 4 18. Linear dynamic range (LDR) of the photodetectors measured at 1 V. The total LDR is 67 dB, comparable to InGaAs, and sufficient for high contrast imaging.

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132 Figure 4 19. ambient conditions without encapsulation. Details are discussed in the text.

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133 CHAPTER 5 CONCLUSIONS AND FUTURE WORK 5.1 Summary This dissertation aimed to spearhead development in the area of nanostructured materials in next generation technologies like organic solar cells and infrared and multispectral photodetectors. Research and development in this field is key to moving technology and our modern society forward, as it can unlock energy inde pendence and inspire new tools for imaging, security, and telecommunications. In chapter 2, we characterized the physical, chemical, optical, and electronic properties of solution processed NiO and fabricated solar cells using this NiO film as a ho le transport layer. Solar cells incorporating NiO films showed significant enhancements in fill factor and short circuit current, leading to a 14.7% increase in PCE compared with the cells with PEDOT:PSS. The enhancements in the NiO devices were due to improv ed optical resonance, nanoscale active layer morphology, increased shunt resistance, and lower series resistance for charge extraction in the NiO devices. This is a significant contribution to the scientific literature and represents a step past the status quo in organic photovoltaic cells. However, processing temperatures for this NiO layer are still too high for practical application with plastic substrates like PET that may dominate the roll to roll printing landscape. Additionally, the need for a low te mperature the market. In chapter 3 we developed a novel low temperature route to forming NiO films and fabricated high efficiency solar cells from these lay ers. The formation of these NiO films took place in a two step process, whereby precursor films were heated to temperatures low enough that PET substrates could withstand the heating, then oxidized by an ozone treatment. The final

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134 NiO films were amorphous and contained a strong NiOOH dipole layer which aids in current extraction and electron blocking in solar cells. Finally, as a proof of concept, solar cells with this new low temperature processed NiO were fabricated on PET substrates. The devices showed a n average PCE of 3.7%, which is comparable to standard small area P3HT:PCBM solar cells made as a community standard on glass. The feasibility of creating these solar cells on PET substrates is promising for future development of this technology as the pho tovoltaics community looks for ways to scale up small lab scale devices into intermediate scale or even large scale commercial printing facilities. In chapter 4, we used the nanostructured materials NiO, ZnO and PbS QDs to fabricate a new class of multispe ctral photodetector one with all solution processed inorganic materials. The devices have extremely low noise, a large linear dynamic range of 67 dB, bandwidth of 36 kHz, and specific detectivities over 1x10 12 cmHz 1/2 /W, higher than commercially availabl e germanium photodiodes and comparable to silicon and InGaAs photodiodes. NiO and ZnO were effective charge blocking layers and provided low dark current that the allowed the low noise levels. Additoinally, the use of these oxide blocking layers as native encapsulating layers sandwiching the PbS created an extremely air stable structure as device performance remained unchanged for five months with storage in a ir. 5.2 Outlook and Future Studies The field of organic photovoltaics is a promising one. In order to realize highly efficient devices, chemists, materials scientists, and device physicists will need to collaborate closely to address the requirements for developing next generation materials with even higher efficiencies that are competitive with Silicon and III V materials. The economic advantage that organics have over materials like Si and III V semiconductors is the ease of scalability through roll to roll processing. An aesthetic and utilitarian advantage organic solar cells have over the current

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135 sil icon and III V and even CdTe solar cells is their flexibility, relative thinness and light weight. These factors should drive architectural concepts to incorporate these devices into things like energy harvesting window screens and paints. Smaller scale ni che applications like solar bags and handheld device chargers have already been demonstrated, and should continue to develop with increasing efficiency of the solar cells. The two major drawbacks of organic photovoltaics at the present time are the lack o f air, moisture and UV light stability and also the low power conversion efficiency in module scale solar cells. Much attention needs to be directed toward solving these two problems, the stability by chemists and materials scientists, and the scale up eff orts by everyone working in the field. There are several different polymer systems capable of power conversion efficiencies over 8% in small lab scale devices, yet no polymer photovoltaic cells have been fabricated on a large scale with those high efficien cies to date. These engineering challenges are those of this generation of researchers. If OPV is to be competitive with, or even beat, silicon on a dollar per watt and lifetime basis, addressing these issues is critical. With only a small portion of the O PV research community publishing work on these matters, the field is wide open for innovative solutions and advancements in scale up and chemical design. Future studies on the materials in this dissertation should elucidate more detailed device physics and chemistry. The role of interface recombination and vertical phase separation in BHJ solar cells is an important one. The difference in recombination types and rates with changing HTLs should correlate to vertical phase separation and donor acc eptor morpho logy and HTL surface energy. Additionally, determining the exact chemical route to the formation of NiO from these solution precursors would help engineer better starting materials. A full long term study would include the mass spectrometry included in thi s dissertation, with the addition of single

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136 crystal X ray crystallography to obtain the solid state structure of these precursor ions. Th ermal desorption spectroscopy/temperature programmed desorption measurements would help elucidate the exact mechanism o f NiO formation and determine the order in which the molecules in the precursor decompose and at what energies they break apart. In the multispectral detectors, a slight excess shot noise was measured and is attributed to a small gain mechanism. It is know n that the dependence of the shot noise on DC bias and current can provide information about the gain mechanisms [135] Device level modeling combined with extensive experimental work on the noise measurements for these detectors could produce interesting data on gain mechanisms commonly seen in quantum dot photodetectors.

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137 APPENDIX A ORGANIC MOLECULAR ST RUCTURES 1. PEDOT:PSS 2. pDTG TPD 3. Oleic acid 4. 1,3 Benzenedithiol 5. PC 70 BM

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138 APPENDIX B LIST OF PATENTS, PUBLICATIONS AND PRESENTATIONS Patents 1. Stable Infrared Photodetectors from Solution J.R. Manders S. Chen, T. H. Lai, E.D. Klump, S. W. Tsang, F. So. U.S. Patent Application Number 61/871,579. Filed August 29, 2013. 2. Processed Ultraviolet Light Detector Based on P N Junctions of Metal J.R. Manders D.Y. Kim, J. Ryu, J.W. Lee, F. So. U.S. Patent Application Number: 61/722,403. Filed November 5, 2012. Peer Reviewed Publications 1. Noise, Air Stable Multispectral Photodetectors Made fr om All Solution Processed J.R. Manders T. H. Lai, Y. An, W. Xu, G. Bosman, F. So, Manuscript submitted for publication 2. Stable, Solution Processed Oxide p D.Y. Kim, J. Ryu, J.R. Manders F. So, ACS Applied Materials and Interfaces Accepted for Publication Published online January 9, 2014 DOI: 10.1021/am4050019. 3. H. Lai, S. W. Tsang, J.R. Manders S. Chen, F. So. Materials Today 16, 11, 424 432 (2013). DOI: 10.1016/j.mattod.2013.10.001 4. Induced Loss Mechanisms in Polymer Inorganic Planar Heterojunction Solar Y. Hsu, J.R. Manders K. Schanze, F. So, ACS Applied Materials and Interfaces 5, 15, 7215 7218 (2013) DOI: 10.1021/am4015606. 5. Processed Nickel Oxide Hole Transport Layers in High Efficiency Polymer J.R. Manders S. W. Tsang, M.J. Hartel, T. H. Lai, S. Chen, C.M. Amb, J.R. Reynolds, F. So, Advanced Functional Materials 23, 23, 2993 3001 (2013). DOI: 10.1002/adfm.201202269 6. J.R. Manders S. W. Tsang, F. So, Journal of Materials Chemistry 22, 46, 24202 24212 (2012). DOI: 10.1039/c2jm33838f 7. Small, C.M. Amb, J. Subbiah, T. H. Lai, S. W. Tsang, J.R. Manders J.R. Reynolds, F. So, Advanced Energy Materials 2, 11, 1333 1337 (2012). DOI: 10.1002/aenm.201200184

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139 Conference Presentations 1. Stable Infrared Through Visible Photodetectors from All Solution Processed Inorganic Semiconductors with Low Noise and High Dete J.R. Manders T. H. Lai, Y. An, W. Xu, J.W. Lee, D.Y. Kim, G. Bosman, F. So. (Oral) NanoFlorida 2013. Gainesville, FL. September 2013. 2. J.R. Manders T. H. Lai, J.W. Lee, F. So. Invited Paper (Oral) SPIE Optics + Photonics 2013. San Diego, CA. August 2013. 3. Processed Nickel Oxide Hole J.R. Manders S. W. Tsang, M.J. Hartel, T. H. Lai, S. Chen, C.M. Amb, J.R. Reynolds, F. So. Inaugural SEC Symposium. Atlanta, GA. February 2013 4. World Lecture Competition 2 nd Place J.R. Manders Londo n, United Kingdom. July 2012. 5. Processed Nickel Oxide Hole Transport Layers in High Efficiency Organic J.R. Manders S. W. Tsang, M.J. Hartel, T. H. Lai, S. Chen, C.M. Amb, J.R. Reynolds, F. So. International Conference on Syntheti c Metals (ICSM), Atlanta, GA. July 2012. 6. to Visible Up H. Lai, D.Y. Kim, J.W. Lee, D.W. Song, J. Ryu, J.R. Manders F. So. International Conference on Synthetic Metals (ICSM), Atlanta, GA. July 2012. 7. and Dithienogermole Based Donor Acceptor Conjugated Polymers for W. Tsang, S. Chen, J.R. Manders C. Small, I. Constantinou, F. So, J.R. Reynolds. International Conference on Synthetic Metals (ICSM), Atlanta, GA. July 2012. 8. Processed Nickel Oxide Hole Transport Layers in High Efficiency Organic J.R. Manders S. W. Tsang, M.J. Hartel, T. H. Lai, S. Chen, C.M. Amb, J.R. Reynolds, F. So. (Oral) Materi als Research Society Spring Meeting, San Francisco, CA. July 2012 9. K.R. Choudhury, L. Qian, J.R. Manders and F. So. 2010 Materials Research Society (MRS) Fall Meeting, Bosto n, MA. Nov. Dec. 2010.

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140 LIST OF REFERENCES [1] H. Shirakawa, E. J. Louis, A. G. Macdiarmid, C. K. Chiang, A. J. Heeger, J. Chem. Soc. Chem. Commun. 1977 1977 578. [2] J. H. Burroughes, D. D. C. Bradley, A. R. Brown, R. N. Marks, K. Mackay, R. H. Friend, P. L. Burns, A. B. Holmes, Nature 1990 347 539. [3] D. Kearns, M. Calvin, J. Chem. Phys. 1958 29 950. [4] C. W. Tang, S. A. Vanslyke, Appl. Phys. Lett. 1987 51 913 [5] C. W. Tang, Appl. Phys. Lett. 1986 48 183. [6] C. E. Small, S. Chen, J. Subbiah, C. M. Amb, S. Tsang, T. H. Lai, J. R. Reynolds, F. So, Nat. Photonics 2011 6 115. [7] B. Streetman, S. Banerjee, Solid State Electronic Devices Pearson Prentice Hal l, 2006 [8] B. Carsten, J. M. Szarko, H. J. Son, W. Wang, L. Lu, F. He, B. S. Rolczynski, S. J. Lou, L. X. Chen, L. Yu, J. Am. Chem. Soc. 2011 133 20468. [9] E. E. Havings, W. ten Hoeve, H. Wynberg, Polym. Bull. 1992 29 119. [10] S. W. Tsang, S. Chen, F. So, Adv. Mater. 2013 25 2434. [11] S. Chen, Electronic Processes in Polymer Solar Cells, University of Florida, 2012 [12] O. D. Jurchescu, J. Baas, T. T. M. Palstra, Appl. Phys. Lett. 2004 84 3061. [13] A. A. Grinberg, S. Luryi, M. R. Pinto, N. L. Schryer, IEEE Trans. Electron Devices 1989 36 1162. [14] Y. B. Zhu, L. K. Ang, J. Appl. Phys. 2011 110 094514. [15] S. Chen, K. R. Choudhury, J. Subbiah, C. M. Amb, J. R. Reynolds, F. So, Adv. Energy Mater. 2011 1 963. [16] S. Chen, C. E. Small, C. M. Amb, J. Subbiah, T. H. Lai, S. W. Tsang, J. R. Manders, J. R. Reynolds, F. So, Adv. Energy Mater. 2012 2 1333. [17] J. J. M. Halls, C. A. Walsh, N. C. Greenham, E. A. Marseglia, R. H. Friend, S. C. Moratti, A. B. Holmes, Nature 1995 376 498. [18] G. Yu, J. Gao, J. C. Hummelen, F. Wudl, A. J. Heeger, Science (80 ). 1995 270 1789.

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148 BIOGRAPHICAL SKETCH Jesse Manders was born to a loving family in Indianapolis, Indiana and soon moved to the suburbs of Dayton, Ohio, where he spent his childhood. After graduating from Centerville High School, he attended Miami University (Ohio) in Oxford, Ohio, where he fir st majored in Chemistry, and soon switched to Engineering Physics Engineering Physics allowed him to pursue his passion of pure physics while adding applied coursework to the curriculum. Although majoring in the physics department, Jesse kept a keen inter est in chemistry, tutoring students in of Standards and Technolo gy (NIST) in the suburbs of Washington, DC, he knew he would follow his passion to be a scientist and engineer. He graduated from Miami University with Cum Laude honors and was inducted into P hi Beta Kappa S ociety. After spending 23 years of winters in th e cold and under dreary skies, Jesse opted for the Sunshine State and the University of Florida for his Ph.D. studies, joining the highly respected Department of Materials Sc ience and Engineering There he found a motivated and fast paced research group to join under the direction of Professor Franky So. The group works on developing electronic and optoelectronic materials geared towards next generation organic solar cells, organic light emitting diodes (OLEDs), now common in smart phones, organic transisto rs, and quantum dot based infrared photodetectors. Jesse learned a broad spectrum of information and has been a lead author or contributing author on six peer reviewed publications the lead inventor on two U.S. patent applications, and has delivered nine conference presentations, including earning second place in the World Lecture Competition in London, England, which was sponsored by Rolls Royce and CBMM and administered by IOM3, for a lecture to a non

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149 expert audience on solar energy. In the spring of 201 4, Jesse was awarded his Ph.D. from the University of Florida.