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Luminescent Oxide Nanocomposite

Permanent Link: http://ufdc.ufl.edu/UFE0042049/00001

Material Information

Title: Luminescent Oxide Nanocomposite Synthesis, Characterization and Scintillation Application
Physical Description: 1 online resource (141 p.)
Language: english
Creator: Tseng, Teng-Kuan
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2010

Subjects

Subjects / Keywords: bgo, luminescent, nanocrystal, phosphor, scintillator
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: Scintillator crystals have been made with complex single crystal growth methods which frequently result in high costs and small crystal size. This work has investigated other alternative processing methods for the preparation of scintillation nanocomposite in order to develop low-cost processes for larger area and mass production of ceramic scintillation materials. A variety of ceramic scintillator oxide nanocomposites were synthesized using either non-hydrolytic hot solution route at 280 ? or aqueous precipitation methods at low temperature of 90 ? in short reaction time of 30 min to 2 h. In the study, 20 nm flower-like and 10 nm round Gd2O3:Eu3+ nanocrystals were synthesized using non-hydrolytic hot solution method. The effects of reaction procedure and temperature upon nanocrystal shape were explained based on nucleation and growth of oxide nanocrystals. In addition, Gd(OH)3:Eu3+ nanocrystals were synthesized through sol-gel precipitation method and evolved from nano-size spheres into one dimensional rod-like nanocrystals ~800 nm long with a diameter of ~70 nm (aspect ratio: ~11) after growth for 2 h. The addition of polyethylene glycol 8000 (PEG-8000) to the reaction resulted in slower growth along the 001 direction, leading to shorter rod-like nanocrystals with a smaller aspect ratio of ~2. The hexagonal Gd(OH)3:Eu3+ were converted to cubic Gd2O3:Eu3+ while maintaining the nanorod geometry and dimensions after the calcination at 800 ? for 2 h. The photoluminescence (PL) and radioluminescence (RL) emission spectra of these calcined Gd2O3:Eu3+ nanocrystals excited by 280 nm ultra-violet (UV) irradiation and 40 KeV X-ray, respectively, showed a dominant 5D0-7F2 transition at 612 nm from Eu3+. Nanostructures with mono-dispersed 220 nm SiO2 cores capped with a 13 nm Gd2O3:Eu3+ shell were prepared via a urea precipitation method. The QY from SiO2/Gd2O3:Eu3+ core/shell nanoparticles excited at 280 nm decreased from 25.0 % to 16.5 % with addition of a SiO2 shell to form a SiO2/Gd2O3:Eu3+/SiO2 nanostructure. In contrast, the QY increased to 32.2 % from a SiO2/Gd2O3:Eu3+/Gd2O3 nanostructure. Moreover, for a SiO2:Eu3+ core, addition of a Gd2O3 capping layer increased the QY by four times. This enhancement was attributed to the crystalline Gd2O3 layer acting as an antenna for energy transfer to the Eu3+. Three dimensional (3D) self-assembled hierarchical Bi2O3 architectures were prepared via a solution precipitation synthesis at 85 ? in 45 min with the aid of PEG-8000 as a capping agent. With an increase concentration of PEG, the morphology and structural phase evolved from monoclinic ?-phase micro-rods into cubic ?-phase flower-like crystals. For flower-like Bi2O3 crystals, the morphology change from 60 nm nano-spheres to agglomerated sub-micron clusters to predominantly 3D self-assembled crystals. Further deposition of Gd2O3:Eu3+ on this hierarchical Bi2O3 led to a luminescent core/shell composite. Intense red light emission was observed under UV irradiation after calcining at 500 ? for 2 h in air. Self-activated Bi4Ge3O12 (BGO) crystals were synthesized at a low reaction temperature ( < 90 ?), in short reaction times ( < 30 min) in a water-based solution. Flower- and coral-like BGO crystals obtained by using different precipitation agents had QYs of 28% and 80%, respectively, when excited at 280 nm after calcination at 600 ? for 2 h. Coral-like BGO crystals after being calcined showed a scintillation response when excited by 241Am and 137Cs irradiation sources. Moreover, Bi2O3/BGO core/shell scintillation composites were synthesized with a one-pot solution precipitation method. The BGO shell nucleated heterogeneously as islands on the Bi2O3 microrod cores, and was followed by diffusion-limited lateral growth into a continuous dendrite-like shell. A broad peak from the 59 KeV ? emission from 241Am was observed in the differential pulse height spectrum from calcined core/shell composite. When excited at 280 nm, these self-activated BGO crystals had a broad photoluminescent emission band from 350 to 700 nm with the peak at 530 nm assigned to the 3P1?1S0 transition of Bi3+. Self-assembled almond-shape colloidal GdVO4:Eu3+ nanocrystals composed of ~60 nm long and ~10 nm wide nanorods were synthesized in aqueous solution. The as-prepared nanocrystals were crystallized in short growth time (~3 min) without further post-grown heat treatment. The PL emission spectra showed strong, sharp peaks near 617 nm associated with the characteristic 5D0-7F2 transitions from Eu3+. For excitation at 347 nm and with 2 mol% co-doped Bi3+ and 10 mol% Eu3+ concentration, the QY was 2.5 times more than that from GdVO4:Eu3+ nanocrystals. This enhancement is attributed to additional absorption from the Bi-O charge transfer band.
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.
Statement of Responsibility: by Teng-Kuan Tseng.
Thesis: Thesis (Ph.D.)--University of Florida, 2010.
Local: Adviser: Holloway, Paul H.
Local: Co-adviser: Davidson, Mark R.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2010
System ID: UFE0042049:00001

Permanent Link: http://ufdc.ufl.edu/UFE0042049/00001

Material Information

Title: Luminescent Oxide Nanocomposite Synthesis, Characterization and Scintillation Application
Physical Description: 1 online resource (141 p.)
Language: english
Creator: Tseng, Teng-Kuan
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2010

Subjects

Subjects / Keywords: bgo, luminescent, nanocrystal, phosphor, scintillator
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: Scintillator crystals have been made with complex single crystal growth methods which frequently result in high costs and small crystal size. This work has investigated other alternative processing methods for the preparation of scintillation nanocomposite in order to develop low-cost processes for larger area and mass production of ceramic scintillation materials. A variety of ceramic scintillator oxide nanocomposites were synthesized using either non-hydrolytic hot solution route at 280 ? or aqueous precipitation methods at low temperature of 90 ? in short reaction time of 30 min to 2 h. In the study, 20 nm flower-like and 10 nm round Gd2O3:Eu3+ nanocrystals were synthesized using non-hydrolytic hot solution method. The effects of reaction procedure and temperature upon nanocrystal shape were explained based on nucleation and growth of oxide nanocrystals. In addition, Gd(OH)3:Eu3+ nanocrystals were synthesized through sol-gel precipitation method and evolved from nano-size spheres into one dimensional rod-like nanocrystals ~800 nm long with a diameter of ~70 nm (aspect ratio: ~11) after growth for 2 h. The addition of polyethylene glycol 8000 (PEG-8000) to the reaction resulted in slower growth along the 001 direction, leading to shorter rod-like nanocrystals with a smaller aspect ratio of ~2. The hexagonal Gd(OH)3:Eu3+ were converted to cubic Gd2O3:Eu3+ while maintaining the nanorod geometry and dimensions after the calcination at 800 ? for 2 h. The photoluminescence (PL) and radioluminescence (RL) emission spectra of these calcined Gd2O3:Eu3+ nanocrystals excited by 280 nm ultra-violet (UV) irradiation and 40 KeV X-ray, respectively, showed a dominant 5D0-7F2 transition at 612 nm from Eu3+. Nanostructures with mono-dispersed 220 nm SiO2 cores capped with a 13 nm Gd2O3:Eu3+ shell were prepared via a urea precipitation method. The QY from SiO2/Gd2O3:Eu3+ core/shell nanoparticles excited at 280 nm decreased from 25.0 % to 16.5 % with addition of a SiO2 shell to form a SiO2/Gd2O3:Eu3+/SiO2 nanostructure. In contrast, the QY increased to 32.2 % from a SiO2/Gd2O3:Eu3+/Gd2O3 nanostructure. Moreover, for a SiO2:Eu3+ core, addition of a Gd2O3 capping layer increased the QY by four times. This enhancement was attributed to the crystalline Gd2O3 layer acting as an antenna for energy transfer to the Eu3+. Three dimensional (3D) self-assembled hierarchical Bi2O3 architectures were prepared via a solution precipitation synthesis at 85 ? in 45 min with the aid of PEG-8000 as a capping agent. With an increase concentration of PEG, the morphology and structural phase evolved from monoclinic ?-phase micro-rods into cubic ?-phase flower-like crystals. For flower-like Bi2O3 crystals, the morphology change from 60 nm nano-spheres to agglomerated sub-micron clusters to predominantly 3D self-assembled crystals. Further deposition of Gd2O3:Eu3+ on this hierarchical Bi2O3 led to a luminescent core/shell composite. Intense red light emission was observed under UV irradiation after calcining at 500 ? for 2 h in air. Self-activated Bi4Ge3O12 (BGO) crystals were synthesized at a low reaction temperature ( < 90 ?), in short reaction times ( < 30 min) in a water-based solution. Flower- and coral-like BGO crystals obtained by using different precipitation agents had QYs of 28% and 80%, respectively, when excited at 280 nm after calcination at 600 ? for 2 h. Coral-like BGO crystals after being calcined showed a scintillation response when excited by 241Am and 137Cs irradiation sources. Moreover, Bi2O3/BGO core/shell scintillation composites were synthesized with a one-pot solution precipitation method. The BGO shell nucleated heterogeneously as islands on the Bi2O3 microrod cores, and was followed by diffusion-limited lateral growth into a continuous dendrite-like shell. A broad peak from the 59 KeV ? emission from 241Am was observed in the differential pulse height spectrum from calcined core/shell composite. When excited at 280 nm, these self-activated BGO crystals had a broad photoluminescent emission band from 350 to 700 nm with the peak at 530 nm assigned to the 3P1?1S0 transition of Bi3+. Self-assembled almond-shape colloidal GdVO4:Eu3+ nanocrystals composed of ~60 nm long and ~10 nm wide nanorods were synthesized in aqueous solution. The as-prepared nanocrystals were crystallized in short growth time (~3 min) without further post-grown heat treatment. The PL emission spectra showed strong, sharp peaks near 617 nm associated with the characteristic 5D0-7F2 transitions from Eu3+. For excitation at 347 nm and with 2 mol% co-doped Bi3+ and 10 mol% Eu3+ concentration, the QY was 2.5 times more than that from GdVO4:Eu3+ nanocrystals. This enhancement is attributed to additional absorption from the Bi-O charge transfer band.
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.
Statement of Responsibility: by Teng-Kuan Tseng.
Thesis: Thesis (Ph.D.)--University of Florida, 2010.
Local: Adviser: Holloway, Paul H.
Local: Co-adviser: Davidson, Mark R.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2010
System ID: UFE0042049:00001


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1 LUMINESCENT OXIDE NANOC OMPOSITE: SYNTHESIS, CHARACTERIZATION AND SCINTILLATION APPLICATION By Teng-Kuan Tseng A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLOR IDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2010

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2 2010 Teng-Kuan Tseng

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3 To my parents, my wife, my son and daughter

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4 ACKNOWLEDGMENTS It has been a long journey in pursuit of the P h.D. degree at the University of Florida. The past years have been full with all kinds of tast e in life. The work presented in this dissertation would not have been possible without the help and support from many people. It has been a blessing and privil ege to work with Dr. Paul H. Holloway. He let me have the chance to learn and experience the charisma of a scholar as a supervisor should. His insurmountable patience and unconditional support help me through the low tides during the research period; his professional attitude and ope n-mind character have influenced me greatly on my perception of the scientific research issues I would also like to thank my supervisory committee: Dr. Mark Davidson, Dr. Franky So, Dr Valentin Craciun and Dr. Timothy Anderson for their advice and support during my work. I would like to acknowledge the hospitality and support from Ludie Harmon in Dr. Holloways group. Moreover, I would like to thank previous and current members in our group for all their assistance and useful discussions, particularly Deba sis Bera, Lei Qian, Evan Law, Jason Rowland, Chris Gorrie, Ricardo Torres, Se rgey Maslov and Jihun Choi. I would like to also thank Dr. Luiz Jacobsohn, Dr. Jiange ng Xue, Ying Zheng, Sang-Hyun Eom, William Hammond and Jason Mayer for the support. In add ition, I would like express appreciation to the staff of MAIC and PERC for their assistance, es pecially Kerry Siebein for her support with TEM analysis. All the unconditional love, trust and support from my parents have strengthened me throughout my study. I would like to share the joy of this accomplishment with my dear parents and sisters. In addition, I am co mpletely grateful to my wife, Hui-Wen, for her consistent support and encouragement through this long journey and for her meticulous care of both myself and our

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5 son, Derek, and daughter, Allyson. Thanks to my kids for motivating me to continue and complete my research work. I love you all.

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6 TABLE OF CONTENTS page ACKNOWLEDGMENTS...............................................................................................................4 LIST OF TABLES................................................................................................................. ..........9 LIST OF FIGURES.......................................................................................................................10 ABSTRACT...................................................................................................................................17 CHAP TER 1 INTRODUCTION..................................................................................................................20 1.1 Background and Motivation .............................................................................................20 1.2 Goals and Objectives........................................................................................................20 1.3 Technical Approach......................................................................................................... .21 1.4 Structure of the Dissertation.............................................................................................21 2 LITERATURE REVIEW.......................................................................................................23 2.1 Fundamental of Colloidal Nanocrystals...........................................................................23 2.1.1 Synthetic Processes for Colloidal Nanocrystals................................................... 23 2.1.2 Surface Physical and Chemistr y Properties of Nanocrystals ............................... 27 2.1.3 Growth Mechanism and Shape Control...............................................................30 2.2 Photoluminescence Enhancement of Nanocrystals.......................................................... 35 2.2.1 Surface Passivation Effect.................................................................................... 35 2.2.2 Sensitization for Energy Transfer Effect..............................................................38 2.3 Inorganic Luminescent Materials..................................................................................... 40 2.3.1 Fundamental of Luminescent Materials ............................................................... 40 2.3.2 Application of Luminescent Materials................................................................. 43 2.4 Inorganic Scintillation Materials......................................................................................45 2.4.1 Development and History of I norganic Scintillation Materials ...........................46 2.4.2 Fundamentals and Criteria of Scintillation Materials .......................................... 48 2.4.3 Synthesis Processes and Applica tion of Scintillation Materials ..........................53 3 SYNTHESIS AND CHARACTERIZATION OF LUMINESCENT ZERO AND ONE DIMENSIONAL GADOLINIUM OXIDE NANOCRYSTALS ........................................... 56 3.1 Introduction............................................................................................................... ........56 3.2 Experimental............................................................................................................... ......57 3.2.1 Non-Hydrolytic Hot Soluti on Phase Synthesis Method.......................................57 3.2.2 Aqueous Sol-Gel Precip itation Synthesis Method ............................................... 58 3.2.3 Characterization...................................................................................................58 3.3 Results and Discussion..................................................................................................... 59 3.3.1 Hot Solution Synthesis of Zero Dimensional Gd2O3:Eu3+ Nanocrystals............. 59 3.3.1.1 Shape control of Gd2O3:Eu3+ nanocrystals........................................... 59

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7 3.3.1.2 Structural and luminescent properties of Gd2O3:Eu3+ nanocrystals......61 3.3.2 Sol-Gel Precipitation Synthesis of One Dimensional Gd2O3:Eu3+ Rods.............64 3.3.2.1 Synthesis of Gd2O3:Eu3+ micro-rods..................................................... 64 3.3.2.2 Capping agent effect............................................................................. 65 3.3.2.3 Structural analysis.................................................................................66 3.3.2.4 Formation mechanism........................................................................... 68 3.3.2.5 Luminescent and scinti llation properties of Gd2O3:Eu3+ nanocrystals........................................................................................ 69 3.4 Conclusions.......................................................................................................................72 4 SYNTHESIS AND LUMINESCENT CH ARACTERISTIC OF EUROPIUM DOP ANTS IN SILICA/GADOLINIUM OXIDE CORE/SHELL SCINTILLATION NANOPARTICLES...............................................................................................................74 4.1 Introduction............................................................................................................... ........74 4.2 Experimental............................................................................................................... ......75 4.2.1 Synthesis of Core/Shell and Core/Multi-Shell Nanoparticles.............................. 75 4.2.2 Characterization...................................................................................................76 4.3 Results and Discussion..................................................................................................... 76 4.3.1 Structural Analysis...............................................................................................76 4.3.2 Thin Film Quantum Yield Measurement............................................................. 80 4.3.3 Luminescent Properties of Core/Multi-Shell Nanoparticles ................................ 81 4.3.4 Luminescent Properties of SiO2:Eu3+ and SiO2:Eu3+/Gd2O3 Nanoparticles........ 82 4.3.5 Photoluminescent Enha ncem ent Mechanism....................................................... 84 4.4 Conclusions.......................................................................................................................85 5 CORE/SHELL COMPOSITE OF SELF-A SSEM BLED HIERARCHICAL BISMUTH OXIDE/EUROPIUM-DOPED GADOLINIUM OXIDE....................................................... 86 5.1 Introduction............................................................................................................... ........86 5.2 Experimental............................................................................................................... ......87 5.2.1 Synthesis of Rod-Like Bi2O3 and Hierarchical Flower-Like Bi2O3.....................87 5.2.2 Synthesis of Bi2O3/Gd2O3:Eu3+ Core/Shell Composite....................................... 87 5.2.3 Characterization...................................................................................................88 5.3 Results and Discussion..................................................................................................... 88 5.3.1 Rod-Like and Hierarchical Flower-Like Bi2O3....................................................88 5.3.1.1 Capping agent effect-cry stal and m orphology study............................ 88 5.3.1.2 Time-dependent growth analysis of hierarchical flower-like Bi2O3.....90 5.3.1.3 Structural and photoluminescent pr operties of hierar chical flowerlike Bi2O3............................................................................................93 5.3.1.4 Possible formation mechanism............................................................. 94 5.3.2 Bi2O3/Gd2O3:Eu3+ Core/Shell Composite............................................................ 95 5.3.2.1 Structural analysis.................................................................................95 5.3.2.2 Photoluminescence analysis.................................................................. 96 5.4 Conclusions.......................................................................................................................97

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8 6 SYNTHESIS AND CHARACTERIZATION OF SELF-ACTIVATED BISMUTH GERMANIUM OXIDE SCINTILLATION MATERIALS ................................................... 98 6.1 Introduction............................................................................................................... ........98 6.2 Experimental............................................................................................................... ......99 6.2.1 Materials...............................................................................................................99 6.2.2 Characterization...................................................................................................99 6.3 Results and Discussion................................................................................................... 100 6.3.1 Flower-Like Self-Activated BGO Crystals........................................................ 100 6.3.1.1 Time-dependent growth analysis........................................................ 100 6.3.1.2 Photoluminescence analysis................................................................ 103 6.3.2 Coral-Like Self-Activated BGO Crystals.......................................................... 105 6.3.2.1 Time-dependent growth analysis........................................................ 105 6.3.2.2 Photoluminescence analysis................................................................ 107 6.3.2.3 Radiation test....................................................................................... 109 6.3.3 Bi2O3/BGO Core/Shell Composites................................................................... 110 6.3.3.1 Time-dependent growth analysis........................................................ 110 6.3.3.2 Formation mechanism of dendrite-like shell...................................... 112 6.3.3.3 Photoluminescent analysis.................................................................. 113 6.3.3.4 Radiation test....................................................................................... 114 6.4 Conclusions.....................................................................................................................115 7 ENHANCED PHOTOLUMINESCENCE OF COLLOIDAL SELF-ASSEMBLED ALMOND-SHAPE D GADOLINIUM VANADIUM OXIDE NANOCRYSTALS BY CO-DOPING EUROPIUM AND BISMUTH...................................................................... 117 7.1 Introduction............................................................................................................... ......117 7.2 Experimental............................................................................................................... ....118 7.2.1 Materials.............................................................................................................118 7.2.2 Characterization.................................................................................................118 7.3 Results and Discussion................................................................................................... 119 7.3.1 Structure and Morphology of GdVO4:Eu3+/Bi3+ Nanocrystals..........................119 7.3.2 Photoluminescent Properties of GdVO4:Eu3+/Bi3+ Nanocrystals....................... 122 7.4 Conclusions.....................................................................................................................125 8 CONCLUSIONS AND FUTURE WORK ........................................................................... 127 8.1 Luminescent Zero and One Dimensional Gd2O3:Eu3+ Nanocrystals............................. 127 8.2 Eu3+ Doped SiO2/Gd2O3 Core/Shell Scintill ation Nanoparticles...................................127 8.3 Self-Assembled Hierarchical Bi2O3/Gd2O3:Eu3+............................................................128 8.4 Self-Activated Bi4Ge3O12 Scintillator Materials............................................................128 8.5 Enhanced Photoluminescence of Self-Assembled Gd2O3:Eu3+, Bi3+ Nanocrystals....... 129 8.6 Future Work................................................................................................................ ....130 LIST OF REFERENCES.............................................................................................................131 BIOGRAPHICAL SKETCH.......................................................................................................141

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9 LIST OF TABLES Table page 2-1 Type of luminescence generate d by different excitation sources. .....................................44 2-2 Selected scintillator materials a nd their radiative tr ansitions [91, 92]. .............................. 51 2-3 The desired properties of an ideal scintillator. ................................................................... 52 3-1 Experimental parameters of flowerand sphere-like Gd2O3:Eu3+ nanocrystals................ 61 4-1 Quantum yields of core/shell and core/s hell/shell nanoparticles excited at 280 nm ......... 82 6-1 Activation energies reported in the literature for selected compounds of interest in our one-pot synthesis [173-175]. ..................................................................................... 113

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10 LIST OF FIGURES Figure page 2-1 SEM micrographs of (a) one-dimensional Gd2O3 nanotubes and (b) spherical and hollow structure Gd2O3 nanoparticles [1, 2]...................................................................... 24 2-2 Transmission electron micrographs (top), opt ical spectra (bo ttom left) and images of aqueous solution of various aspect ratios of gold colloids (bottom right). Aspect ratio: (a) 1.35, (b) 1.95, (c) 3.06, (d ) 3.50 and (e) 4.42 [10].......................................................25 2-3 The experimental flow chart of the preparation of CdS:Mn/ZnS core/shell nanocrystals synthesized via a reverse m icelle method with the in set illustrating the structure of reverse micelle [14]........................................................................................ 26 2-4 Size-tunable photoluminescence of CdSe nanocrystals synthesized by hot solution injection m ethod [18].........................................................................................................27 2-5 Illustration of evolution of energy bandgap for bulk crysta ls and nanocrystals [18]. ....... 28 2-6 Illustration of core/shell structure of nanocrystals. (a) Surface m odified core materials with organic molecules. (b) Complete shell coverage of core materials. (c) Core material with a thin shell layer. (d ) Encapsulation of shell materials into core materials. (e) Hollow core/shell structure. (f) Core/multi-shell structure [33].................. 29 2-7 Illustration of nucleation and growth diagram S=Sc: cr itical saturation of monomer concentration to induce nucleation. S=1: equilibrium monomer concentration below which growth stops [34].....................................................................................................31 2-8 Illustration of the anisotropic growth of CdSe nanocrystals by varying the m onomer concentration [46]............................................................................................................. .33 2-9 Anisotropic growth along [001] as a function of capping surfactants concentration leading to various aspect-ratios of CdSe nanorods. (a) short nanorods (aspect ratio:5.5), (b) m edium nanorods (aspect ratio:11), and (c) l ong nanorods (aspect ratio:20) [51]......................................................................................................................34 2-10 Anisotropic disk-shaped nanocrystals of (a) Co and (b) CuS grew along [100] and [110] directions due to selective bonding of the capping surfactant on [001] facets [52-55]................................................................................................................................34 2-11 (a) Schematic illustration of the self -assem bly of rare-earth oxide nanoplates. Examples of the self assembled nanoplates of and the examples on (b) Eu2O3 (c) ZnO, and (d) Nb2O3 [56]....................................................................................................35

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11 2-12 TEM images illustrating the dependence of the shape of Gd2O3:Eu3+ nanocrystals on the Gd precursor: (a) Gd-ace toacetonate, (b) Gd-acetate, and (c) Gd-chloride [57, 58]. ............................................................................................................................................35 2-13 Schematic diagram of inorganic and organic capping shell on a lum inescent nanocrystal core [15].......................................................................................................... 36 2-14 The effects of organic and inorganic she lls on the PL spectra and QYs fr om CdS:Mn luminescent nanocrystal cores [15]....................................................................................37 2-15 Schematic illustration of energy transfer by the exchange interaction [65]. ..................... 39 2-16 Schematic diagram of luminescence from (a) fluorescen ce, denoted as (f) and (b) phosphorescence, denoted as (p) [83]................................................................................41 2-17 (a) Stokes shift is the energy differe nce between absorption and photolum inescence emission. (b) Radiative recombination proce sses in extrinsic luminescence [14]............. 43 2-18 Schematic diagram of a mercury-discharge fluorescent lamp illustrating the light generation procedure and possi ble em ission colors [89]................................................... 45 2-19 (a) Schematic diagram of blue LED with YAG:Ce3+ phosphor coating. (b) Emission spectrum illustrating a combination of blue LED and YAG:Ce3+ phosphor coating [90].....................................................................................................................................45 2-20 Timeline for the discovery of important inorganic scintillator m aterials [96]................... 47 2-21 Schematic diagram of a typical scintillato r detector. .........................................................48 2-22 (a) Differential pulse height distribution for gamm a rays from 137Cs source. (b) The scintillation pulse he ight spectrum from 55Fe measured with a YAP:Ce scintillator at room temperature............................................................................................................... 49 2-23 Illustration of the mechanisms involved in the scintillation pro cess in an inorganic m aterial [96].......................................................................................................................50 2-24 (a) Schematic diagram of computed tomography (CT) imaging. (b) A CT scan image of a transversal slice of hum an body at the level of th e kidneys [99]................................ 54 2-25 (a) Schematic diagram of positron em ission tom ography (PET) with the inset illustrating BGO detectors coupled with PMT. (b) A PET scan image of a brain from Alzheimers disease patient [102, 103]..............................................................................55 3-1 Flow chart of one-pot non-hydrolytic hot solution synthesis procedure. .......................... 58 3-2 High-resolution TEM images of as-prepared Gd2O3:Eu3+ nanocrystals synthesized with different experimental parameters le ading to nanocrystals with shapes of (a) flowerand (b) sphere -like nanocrystals............................................................................ 61

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12 3-3 XRD pattern of sphere-like Gd2O3:Eu3+ nanocrystals....................................................... 62 3-4 (a) Photoluminescence excitation (PLE) spectrum of as-prepared sphere-like Gd2O3:Eu3+ nanocrystals for the emission peak at 609 nm. (a) Photoluminescence (PL) spectrum sphere-like Gd2O3:Eu3+ nanocrystals excited by 260 nm. The insets show red emission from solid state powd er and colloidal solution samples under UV excitation..................................................................................................................... .......63 3-5 Radioluminescence spectrum of as-prepared sphere-like Gd2O3:Eu3+ nanocrystals......... 64 3-6 SEM micrographs of sol-gel Gd(OH)3:Eu3+ at reaction times of (a) 3 min showing nano-size spheres, (b) 30 mi n showing a mixture of sphere and rod-like particles, and (c) micro-rods after 2 h............................................................................................... 65 3-7 SEM micrographs of Gd(OH)3:Eu3+ nano-structures synthesized by sol-gel with PEG capping agent for reaction times of (a) 15 min and (b) 2 h................................................ 66 3-8 XRD patterns of as-prepared and calcined m icro-rods samples grown for 2 h. (a) Asprepared samples; (b) Diffraction peaks from JCPDS card #83-2037 for hexagonal Gd(OH)3; (c) Calcined samples; (d) Diffraction peaks from JCPDS card #43-1014 for cubic Gd2O3..................................................................................................................67 3-9 (a) TEM micrograph of as-prepared Gd(OH)3:Eu3+ micro-rods grown for 2 h. (b) High-resolution TEM image (Scale bar: 5 nm ); the inset is a selected area electron diffraction pattern (SAED)................................................................................................68 3-10 SEM micrographs of 1D micro-rods of (a) as-prepared hexagonal Gd(OH)3:Eu3+ and (b) calcined cubic Gd2O3:Eu3+...........................................................................................68 3-11 Photoluminescence spectra from Gd2O3:Eu3+ micro-rods (black) and nano-rods (red). (a) PLE spectra for emission at 612 nm. (b ) PL spectra excited at 280 nm. Decay time for (c) nano-rods and (d) micro-rods......................................................................... 71 3-12 Radioluminescence spectrum of calcined Gd2O3:Eu3+ micro-rods.................................... 72 4-1 (a) SEM micrograph of mono-dispersed SiO2 nanoparticles synthesized by the Stober method. TEM micrographs of as-prepared SiO2/Gd2O3:Eu3+ core/shell nanoparticles at (b) low magnification and (c) high magnification. (d) TEM micrograph of SiO2/Gd2O3:Eu3+ core/shell nanoparticles after being calcined at 800 for 2 h in air. The inset shows the luminescence under room light or UV irradiation from a Hg discharge lamp................................................................................ 77 4-2 XRD spectra obtained from (a) pure silica nanop articles, (b) as-prepared SiO2/Gd2O3:Eu3+ core/shell nanoparticles, and (c) calcined SiO2/Gd2O3:Eu3+ core/shell nanoparticles (800 for 2 h in air). (d) Standard diffraction peaks from JCPDS (#43-1014) for cubic Gd2O3..................................................................................78

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13 4-3 TEM micrographs of core/shell/she ll nanostructure of (a) as-prepared SiO2/Gd2O3:Eu3+/SiO2 nanoparticles and (b) calcined SiO2/Gd2O3:Eu3+/SiO2 nanoparticles......................................................................................................................79 4-4 (a) Optical transmission spectra obtai ned from as-prepared and calcined SiO2 nanoparticle layer. (b) XRD spectra of as-prepared and calcined SiO2 nanoparticle layer....................................................................................................................................80 4-5 Schematic diagram of thin film quantum yield (QY) measurement.................................. 81 4-6 Photoluminescence spectra from SiO2/Gd2O3:Eu3+ core/shell and SiO2/Gd2O3:Eu3+/SiO2 core/shell/shell nanoparticles. (a) PLE spectra for emission at 609 nm. (b) PL spectra excited at 280 nm......................................................................... 82 4-7 (a) Photoluminescence spectra from SiO2:Eu3+ and SiO2:Eu3+/Gd2O3 core/shell nanoparticles with PLE spectra for emission at 609 nm and PL spectra excited at 278 nm. (b) Quantum yield of SiO2:Eu3+ and SiO2:Eu3+/Gd2O3 as a function of Gd2O3 reaction time with the inse ts showing photos of PL excited by UV irradiation................ 83 4-8 Schematic illustration of possibl e charge transfer paths in S iO2:Eu3+/Gd2O3 core/shell and SiO2/Gd2O3:Eu3+/Gd2O3 nanoparticles....................................................... 84 5-1 SEM images of morphology evolution as a function of PEG-8000 volum e fraction, C = (a) 0.05, (b) 0.1, and (c) 0.2 (reaction temperature = 85 time = 45 min)................. 89 5-2 X-ray diffraction (XRD) pattern from bis muth oxide as a function of PEG-8000 volume fraction, C = (a) 0.05, (b) 0.1, a nd (c) 0.2 (reaction temperature = 85 time = 45 min). (d) Diffraction peaks from JCPDS for monoclinic -bismuth oxide. (e) Diffraction peaks from JCPDS for cubic -bismuth oxide................................................ 90 5-3 Evolution of the morphology of bismuth oxide as a function of reaction times of (a) 1 m in, (b) 10 min and (c) 45 min (C = 0.2, reaction temperature = 85 ).......................... 91 5-4 XRD pattern from bismuth oxide as a function of reaction times of (a) 1 min, (b) 10 m in and (c) 45 min (black: as-grown; red: calcined at 600 for 2 h in air), (C = 0.2, reaction temperature = 85 ). (d) Diffraction peaks from JCPDS for cubic -bismuth oxide...................................................................................................................................92 5-5 SEM micrographs at different m agnifications of flower-like Bi2O3 calcined at 600 for 2 h in air showing that calcining does not change the morphology............................. 92 5-6 (a) SEM micrograph of an individua l flower of bismuth oxide. (b) TEM micrograph of the lattice fringes from an individual nano-tria ngle with its SAED pattern shown in the inset.................................................................................................. 93 5-7 Photoluminescence spectrum of as-synthesized flower-like -bism uth oxide for excitation at 230 nm........................................................................................................... 94

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14 5-8 Schematic illustration of a possible form ation process for 3D flower-like bism uth oxide...................................................................................................................................95 5-9 SEM micrographs of (a) self-assemble d flower-like bism uth oxide, (b) bismuth oxide/europium doped gadolinium oxide core/shell scintillation composite and (c) high magnification from the indi vidual core/shell composite........................................... 96 5-10 Photoluminescent analysis from bism uth oxide/gadolinium oxide core/shell scintillation composite. (a) PLE spectra obser ved at 612 nm. (b) PL spectra excited at 280 nm. The inset shows the photos under room lighting and UV irradiation.................. 97 6-1 SEM micrographs showing the morphology a nd particle size of as-prepared flowerlike BGO crystals at reaction tim es of (a) 2 min (b) 35 min and (c) 1 h......................... 100 6-2 SEM micrographs of flower-like BGO crystals after growth for 1 h: as-prepared at (a) low and (b) high m agnification, a nd (c) after calcining at 500 for 2 h in air.............. 101 6-3 X-ray diffraction (XRD) pattern from as -prepared flower-like BGO crystals as a function of reaction tim e: (a) 2 min, (b ) 35 min and (c) 1h. (d) Diffraction peaks from JCPDS for cubic Bi4Ge3O12 (blue pattern)..............................................................102 6-4 X-ray diffraction (XRD) pattern from flower-like BGO crystals calcined at 500 for 2 h in air for reaction times of (a) 2 min, (b) 35 min and (c) 1h. (d) Diffraction peaks from JCPDS for cubic Bi4Ge3O12 (blue pattern) and Bi12GeO20 (red pattern)...... 103 6-5 Photoluminescence spectra from flower-l ike and single BGO crystal sam ples. (a) Photoluminescence excitation (PLE) spectru m for emission at 467 nm from calcined sample (400 2 h). (b) Photoluminescence (PL) spectra of as-prepared and calcined BGO polycrystalline or single B GO crystal samples excited at 280 nm. The insets show color photographs of the emission excited by UV irradiation and the corresponding quantum yields......................................................................................... 104 6-6 SEM micrographs of as-prepa red coral-like BGO crystals at reaction times at (a) 3 min (b) 30 min and (c) 2 h............................................................................................... 105 6-7 (a) SEM micrographs of as-prepared cora l-like BGO sa mples after 2 h of growth. The inset shows individual nanocrystal s at a higher magnification. (b) TEM micrograph of as-prepared coral-like BGO crystals after 2 h growth. (c) Low and (d) high magnification SEM micrographs of co ral-like BGO samples calcined at 600 for 2 h...............................................................................................................................106 6-8 X-ray diffraction (XRD) patterns from as-prepared (black line) and calcined (red line) coral-like BGO structures for reaction tim es of (a) 3 m in, (b) 30 min and (c) 2h. (d) Diffraction peaks from JCPDS for cubic Bi4Ge3O12 (blue pattern) and Bi12GeO20 (orange pattern)............................................................................................................... .107

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15 6-9 Photoluminescence (PL) and photolumin escence excitation (P LE) spectra from asprepared coral-like (a) BGO and (c) BGO:Eu3+ crystals where PLE spectra used emission at 467 nm, and PL spectra were excited at 280 nm. The CIE coordinates from as-prepared coral-like (b) BGO and (d) BGO:Eu3+ crystals excited at 280 nm...... 108 6-10 Differential pulse height spectrum of coral-like BGO crystals excited by (a) 241Am and (b) 137Cs gamma-ray irradiation................................................................................110 6-11 SEM micrographs of Bi2O3/BGO core/shell composite at different reaction times: (a) 30 min, (b) 1 h, (c) 2 h, (d) 3 h, (e) 4 h, and (f) high magnification SEM micrograph of surface dendrite-like feature for a 4 h reacted sample................................................. 111 6-12 SEM micrographs of Bi2O3/BGO core/shell composite at reaction times of (a) 3 h (shell coverage incomplete), and (b) complete shell coverage at 4 h.............................. 111 6-13 X-ray diffraction (XRD) sp ect ra from as-prepared Bi2O3/BGO core/shell composite as a function of reaction time: (a) 30 min, (b) 4 h, (c) 4 h sample calcined at 500 2 h in air, and (d) diffraction peaks from JCPDS files for cubic (blue pattern) and tetragonal (orange pattern) Bi2O3..................................................................................... 112 6-14 (a) Photoluminescence excitation (emission at 530 nm ), and (b) emission (excitation at 280 nm) spectra from calcined Bi2O3/BGO core/shell composite. The inset shows photographs of emission irradiated by room light (top) and UV (bottom)......................114 6-15 Differential pulse hei ght spectrum of calcined Bi2O3/BGO core/shell composite irradiated by an 241Am source showing a broad scin tillation response centered at approximately channel 30................................................................................................ 115 7-1 X-ray diffraction (XRD) patterns from (a ) as-prepared, (reacted for 3 m in. at <90 ), (b) calcined (800 2 h in air) GdVO4:Eu3+ nanocrystals, and (c) diffraction peaks from JCPDS (#86-0996) for cubic GdVO4...........................................................120 7-2 SEM micrographs of as-prepared GdVO4:Eu3+ nanocrystals after reaction times of (a) 3 min and (b) 2 h, and calcined samples (800 2 h in air) reacted for (c) 3 min and (d) 2 h prior to calcining........................................................................................... 120 7-3 TEM micrographs of GdVO4:Eu3+ nanocrystals reacted for 3 min showing almondshaped clusters self-assemb led from nanorods. Note the change from low to high magnification in (a)-(c). (d) High-reso lution TEM image showing well-resolved lattice fringes. Inset: selected area electron diffraction pattern (SAED).........................121 7-4 TEM micrographs, at progressively higher m agnification in (a)-(d), of GdVO4:Eu3+ nanocrystals after 2 h of growth. Note that the self-assembled almond-shaped clusters were of the same size (30~60 nm) as those grown for 3 min............................. 122 7-5 (a) Room temperature photoluminescence excitation (PLE) s pectra from as-prepared GdVO4:Eu3+ (10%) nanocrystals grown for 2 h, a nd with various concentrations of co-doped Bi3+ ( emission = 617 nm). (b) Colored photogr aphs showing the intensity

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16 and color of emission from GdVO4:Eu3+ and co-doped GdVO4:Eu3+, Bi3+ nanocrystals excited by a handheld mercury discharge lamp that emitted ultraviolet (UV) photons over wavelength regions denoted as a short or l ong spectrum [194]........ 123 7-6 Photoluminescence (PL) sp ectra of as-prepared GdVO4:Eu3+ (10%) nanocrystals (grown for 2 h) co-doped with concentrations of Bi3+ ranging from 0-15%, with an excitation wavelength of (a) 280 nm (Inse t: emission from nanocrystals in a quartz cuvette), (b) 323 nm and (c) 347 nm................................................................................ 125

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17 Abstract of Dissertation Pres ented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy LUMINESCENT OXIDE NANOC OMPOSITE: SYNTHESIS, CHARACTERIZATION AND SCINTILLATION APPLICATION By Teng-Kuan Tseng August 2010 Chair: Paul H. Holloway Major: Materials Science and Engineering Scintillator crystals have b een made with complex single crystal growth methods which frequently result in high costs and small crystal si ze. This work has investigated other alternative processing methods for the preparation of scinti llation nanocomposite in order to develop lowcost processes for larger area and mass production of ceramic scintillation ma terials. A variety of ceramic scintillator oxide nanocomposites were synthesized using eith er non-hydrolytic hot solution route at 280 or aqueous precipitation met hods at low temperature of 90 in short reaction time of 30 min to 2 h. In the study, 20 nm flower-like and 10 nm round Gd2O3:Eu3+ nanocrystals were synthesized using non-hydrolytic hot solution method. The effects of reaction procedure and temperature upon nanocrystal shap e were explained based on nuc leation and growth of oxide nanocrystals. In addition, Gd(OH)3:Eu3+ nanocrystals were synthesized through sol-gel precipitation method and evolve d from nano-size spheres into one dimensional rod-like nanocrystals ~800 nm long with a diameter of ~70 nm (aspect ratio: ~11) after growth for 2 h. The addition of polyethylene glycol 8000 (PEG-8000) to the reaction resulted in slower growth along the [001] direction, l eading to shorter rod-like nanocrystals with a sm aller aspect ratio of

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18 ~2. The hexagonal Gd(OH)3:Eu3+ were converted to cubic Gd2O3:Eu3+ while maintaining the nanorod geometry and dimensions after the calcination at 800 for 2 h. The photoluminescence (PL) and radioluminescence (RL) emission spectra of these calcined Gd2O3:Eu3+ nanocrystals excited by 280 nm ultra-violet (UV) irradi ation and 40 KeV X-ray, respectively, showed a dominant 5D0-7F2 transition at 612 nm from Eu3+. Nanostructures with mono-dispersed 220 nm SiO2 cores capped with a 13 nm Gd2O3:Eu3+ shell were prepared via a urea precipitation method. The QY from SiO2/Gd2O3:Eu3+ core/shell nanoparticles excited at 280 nm decreased from 25.0 % to 16.5 % with addition of a SiO2 shell to form a SiO2/Gd2O3:Eu3+/SiO2 nanostructure. In contrast, th e QY increased to 32.2 % from a SiO2/Gd2O3:Eu3+/Gd2O3 nanostructure. Moreover, for a SiO2:Eu3+ core, addition of a Gd2O3 capping layer increased the QY by four times. This enha ncement was attributed to the crystalline Gd2O3 layer acting as an antenna for energy transfer to the Eu3+. Three dimensional (3D) self-assembled hierarchical Bi2O3 architectures were prepared via a solution precipitation synthesis at 85 in 45 min with the aid of PEG-8000 as a capping agent. With an increase concentration of PEG, the morphology and structural phase evolved from monoclinic -phase micro-rods into cubic -phase flower-like crys tals. For flower-like Bi2O3 crystals, the morphology change from 60 nm nanospheres to agglomerated sub-micron clusters to predominantly 3D self-assembled crystals. Further deposition of Gd2O3:Eu3+ on this hierarchical Bi2O3 led to a luminescent core/shell com posite. Intense red light emission was observed under UV irradiation after calcining at 500 for 2 h in air. Self-activated Bi4Ge3O12 (BGO) crystals were synthesized at a low reaction temperature (<90 ), in short reaction times (<30 min) in a water-based solution. Flowerand coral-like BGO crystals obtained by usi ng different precipitation agents had QYs of 28% and 80%,

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19 respectively, when excited at 280 nm after calcin ation at 600 f or 2 h. Coral-like BGO crystals after being calcined showed a scin tillation response when excited by 241Am and 137Cs irradiation sources. Moreover, Bi2O3/BGO core/shell scintillation composite s were synthesized with a onepot solution precipitation method. The BGO shell nucleated heterogeneously as islands on the Bi2O3 microrod cores, and was followed by diffusionlimited lateral growth into a continuous dendrite-like shell. A broad peak from the 59 KeV emission from 241Am was observed in the differential pulse height spectru m from calcined core/shell com posite. When excited at 280 nm, these self-activated BGO crystals had a broa d photoluminescent emission band from 350 to 700 nm with the peak at 530 nm assigned to the 3P11S0 transition of Bi3+. Self-assembled almond-shape colloidal GdVO4:Eu3+ nanocrystals composed of ~60 nm long and ~10 nm wide nanorods were synthe sized in aqueous solution. The as-prepared nanocrystals were crystallized in short growth time (~3 min) without further post-grown heat treatment. The PL emission spectra showed stron g, sharp peaks near 617 nm associated with the characteristic 5D0-7F2 transitions from Eu3+. For excitation at 347 nm and with 2 mol% co-doped Bi3+ and 10 mol% Eu3+ concentration, the QY was 2.5 times more than that from GdVO4:Eu3+ nanocrystals. This enhancement is attributed to additional absorption from the Bi-O charge transfer band.

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20 CHAPTER 1 INTRODUCTION 1.1 Background and Motivation Phosphors are luminescent materials that e fficiently emit light upon electromagnetic or particle excitation. The energy of the excited phot ons/particles is absorbed and dispersed through the materials in the form of energetic electrons and holes. Eventually, these high energy charge carriers relax toward the electronic ground state by either a radiative relaxation resulting in luminescence, or non-radiative relaxation dissipa ting into heat. A scintillator is a phosphor material with the ability to absorb energetic radiation and to convert a fraction of the absorbed energy into photons with characteri stic energies (infra red, visible or ultraviolet photons). Often scintillation materials are modified by incorporat ing dopants (activators) i ons into the host matrix to produce quantum states in the forbidden ba ndgap on which charge carriers are trapped and from which radiative relaxation may occur. Detection of X-ray, -ray, neutrons and charged particle s involves the energy deposited by the absorbed radiation an d generation of secondary hot electron s or charged particles. Inorganic scintillator crystals have traditionally been ma de with complex single crystal growth methods such as Czochralski and Bridgman methods, whic h frequently result in high costs and small crystal size. Therefore, development of low-cost processes for larger area polycrystalline ceramic scintillation materials with high luminescence is of great interest due to their potential for mass production, versatility in shap es and size, and low cost. 1.2 Goals and Objectives The goal of this dissertation is to synthesize scintillator nanocrystals for the utilization in the next generation of ceramic scintillation detect ors. The primary objective of this research was to synthesize a variety of scintillating nanocryst als and to characterize th eir optical properties,

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21 particularly with respect to lu minescence excited by UV, X-ray and -ray photons. It is envisioned (but not proved by this research) that these synt hesis routes for nanocomposite scintillators should be advantag eous in terms of simplicity an d cost versus the techniques required to synthesize single crysta l scintillators. This research did demonstrate that aqueous solution precipitation methods can be used to produce scintillating nanoparticl es at relatively low temperatures (<90 ) in short reaction times (30 min to 2 h). In addition, it was shown that a core/shell structure or co-doping could lead to improved luminescence excited by radioactive sources. 1.3 Technical Approach In order achieve the objective stat ed above, the following was accomplished. 1. Synthesized zeroand one-di mensional scintillation nanocomposites using a facile waterbased solution precipitation method. 2. Characterized scintillation nanocomposites for crystallinity, morphology, structure phase and photoluminescence. 3. Demonstrated that core/shell structure and en ergy transfer from sensitizers enhanced the light output from scintillation nanocomposites. 4. Characterized radioluminescence from the scintillators under excitation by X-ray, 241Am and 137Cs sources. 1.4 Structure of the Dissertation The dissertation research topi c is introduced in Chapter 1. A review of luminescent nanocrystals and their applicati ons in different fiel ds is presented at Chapter 2. Relevant techniques and processes which motivated the de velopment of inorganic luminescent materials were discussed. In addition, strategies to improve photoluminescence were compared and reviewed. Chapter 3 discusses preparation and characterization of zer oand one-dimensional europium-doped gadolinium oxide nanocrystals. Chap ter 4 details the effect s of core/shell and multi-shell structures and dopant location on photoluminescence. A unique self-assembled

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22 hierarchical bismuth oxide cr ystal synthesis by an aqueous solution precipitation method is reported in Chapter 5. Chapter 6 describes th e synthesis and charac terization of bismuth germinate (BGO) crystals prepared at low temperature in short time. Chapter 7 discusses how a sensitizer affects photoluminescence from colloidal self-assembled GdVO4 nanocrystals. The conclusions and future work from this research are summarized in Chapter 8.

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23 CHAPTER 2 LITERATURE REVIEW 2.1 Fundamental of Colloidal Nanocrystals Over the last decade, synthesis and applic ation of inorganic nanocrystals have been intensively studied, not only for their fundamental scie ntific properties, but also for their many potential applications. Nanocrystal s exhibit interesting size-depende nce in their optical, magnetic, chemical, and/or electrical properties which can not be achieved in their bulk counterparts. Several physical and chemical synthetic methods ha ve been used to prepare colloidal, sometimes anisotropic nanocrystals of semiconductor and metal oxide materials. 2.1.1 Synthetic Processes for Colloidal Nanocrystals In general, there are two different routes to synthesize nanocryst als, i.e. top-down approach, which use physical methods, and the bottom-up approach, which utilize solutionphase colloidal chemistry. In our study, we fo cus primarily on the colloidal solution based synthesis methods due to their advantages in uniform nanocrystals with controlled size. In addition, various shapes of nanocrystals can be obtained, e.g. zeroto two-dimensional nanocrystals, by manipulating the reaction cond itions. In the following sections, several synthesis methods wi ll be discussed. Hydrothermal synthesis method has been used in various types of nanocrystals such as TiO2, ZrO2, Ge nanowire and rare-earth doped Gd2O3 [1-6]. The method exploits the precursor solution at elevated temperatures and pressures above its critical poi nt to increase the solubility of reagents and to speed up reac tions between solids. Subsequen tly, precipitation of nanocrystals occurred from the precursor reactions. During the synthesis process, parameters such as reaction pressure, temperature, reacti on time and surfactants can be manipulated to achieve a high simultaneous nucleation rate and multi-dimensi onal morphology and size. The process is called

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24 solvothermal synthesis if othe r non-aqueous polar or nonpolar solv ents are used in the system leading to highly anisotropic nanocrystals. An example is shown in Figure 2-1(a) of onedimensional Gd2O3 nanotubes synthesized in a Teflon autoclave at 180 for 24 h with the pH value of 13. By using different precipitation ag ent and reaction temper ature, spherical and hollow structure Gd2O3 nanoparticles can be obtained via th is hydrothermal synthesis method as shown in Figure 2-1(b). Figure 2-1. SEM micrographs of (a) one-dimensional Gd2O3 nanotubes and (b) spherical and hollow structure Gd2O3 nanoparticles [1, 2]. Reduction synthesis has been reported in the literature to prepare metal and other oxide nanocrystals [7]. In generally, various reagents were used such as sodium borohydride, hydrogen and alcohols. Faraday, et al. prepared colloidal gold nanopartic les from the reduction of HAuCl4 with phosphorous [8]. Later, stable ~10 nm co lloidal gold nanoparticles were synthesized by using sodium citrate as both a reductant and a stabilizer [9]. Sodium borohydride is a strong reducing agent and the reaction takes place immedi ately to generate mono-dispersed colloidal nanoparticles. In contrast, reduction using s odium citrate needs a temperature around 80 Figure 2-2 shows typical colloidal gold nanoparticles with various shapes and sizes synthesized by the reduction reaction.

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25 Figure 2-2. Transmission electron mi crographs (top), optical spectra (bottom left) and images of aqueous solution of various aspect ratios of gold colloids (bottom right). Aspect ratio: (a) 1.35, (b) 1.95, (c) 3.06, (d ) 3.50 and (e) 4.42 [10]. For the reverse micelle synthesis method, the hydrophilic head groups are directed toward the core of the micelles and the hydrophobic groups are directed outward. Reverse micelles are formed in surfactant-stabilized water-in-oil microemulsions and can be formed by ionic surfactants with double long alkyl chains, such as diethyl sulf osuccinate (DES) or a mixture of ionic and nonionic surfactants with a short oxye thylene chain dissolved in organic solvents. Some surfactants like sodium dodecyl sulfate (SDS), cetyltrim ethylammonium bromide (CTAB) are widely used during the reverse micelles sy nthesis. These surfactan ts possess both hydrophilic and hydrophobic groups in the co ntinuous oil phase. This reverse micelle system is heterogeneous on the molecular scale but is thermodynamic stable. Water pools in reverse micelles system serve as reactors in the nanom eter range which favor the formation of small crystalline nanocrystals with narrow size distribution. Continuous dynamic collision with the micellar reactors enables the gr owth of nanocrystals, but the growth beyond the dimension of micellar reactors is inhibited. Metal oxide nanop articles can be prepared inside the reverse micelles by the hydrolysis procedure where metal alkoxide dissolved in oi l reacts with water

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26 inside the droplets. Several metal oxide nanoc rystals have been prepared by this method, including ZrO2, TiO2, SiO2, Fe2O3, silver halide, CaCO3 Y2O3:Eu3+and Gd2O2S:Eu3+ [11-13]. For the synthesis of luminescent semiconductor na nocrystals, Yang, et al. reported synthesis via a reverse micelles method of 2 nm colloidal Cd S:Mn quantum dots passivated by an epitaxial wider band gap ZnS shell to enhance the stabil ity and photoluminescence [14, 15]. The water-tosurfactant ratio for the gr owth of the ZnS shell was critical to have robust and efficient core/shell semiconductor nanocrystals. Figure 2-3 shows the experimental flowchart for the synthesis of efficient CdS:Mn/ZnS core/shell nanocrystals and further characterization procedures. The inset illustrates a schematic cartoon of the nano-size water pool reactor surrounded by surfactant molecules. Figure 2-3. The experimental flow chart of the preparation of CdS:Mn/ZnS core/shell nanocrystals synthesized via a reverse micelle method with the in set illustrating the structure of reverse micelle [14]. The thermal pyrolysis reactions of orga nometallic compounds and metal-surfactant complexes were processed in non-hydrolytic hot surfactant solutions followed by homogeneous nucleation and crystal growth of various semiconductor or metal oxide nanocrystals. The

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27 Bawendi group employed the thermal decomposition method to synthesize mono-dispersed cadmium chalcogenides nanocrystals [16, 17]. Th is method was widely used to synthesize various nanocrystals due to its advantages, such as high crystallinity, mono-dispersed crystals size and their dispersion ability in organic solvents. In a typical synthesis of CdSe nanocrystals, solution of dimethylcadmium ((CH3)2Cd) and tri-n-octylphosphine selenide (TOPSe) were rapidly injected into hot coordinating solvents e.g. tri-n-ocotylphosphine oxide (TOPO) at the temperature of 120-300 inducing a short burst of homogeneous nucleation. The subsequent growth at lower temperature generates CdSe nanocrystals in the range of 1.2 to 12 nm by varying the experimental conditions. These TOPO-capped colloidal CdSe na nocrystals are stable and can be redispersed in organic solvent due to the TO PO surfactant provides a steric passivation of nanocrystal surfaces. The size of nanocrystals sy nthesized by this hot solution injection method is controlled by the reaction temperature and ti me. Due to the quantum confinement effect of CdSe nanocrystals, different crys tal sizes exhibit characteristic emission wavelength as shown in Figure 2-4. Figure 2-4. Size-tunable photolum inescence of CdSe nanocrystals synthesized by hot solution injection method [18]. 2.1.2 Surface Physical and Chemistry Properties of Nanocrystals Nanocrystals are small scale crystals with a size in the range of 1-100 nm. As discussed above, nanocrystals may exhibit novel electric, optical, magne tic, mechanical and chemical properties different from their bulk counterpart materials due to the quantized energy states

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28 and/or the quantum confinement effect [19]. Fi gure 2-5 shows the energy band gap for different nanocrystal sizes and the bulk crys tal. The atomic orbitals of the individual atoms approach each other and generate bonding and anti-bonding molecu lar orbitals. For bulk materials, the energy band structure is continuous with a smaller forbidden bandgap. As the nanocrystal size decreases, the number of atomic orbitals decreases, evolving into an energy band with discrete energy states and a larger forbidden bandgap than bulk material due to constrained electron-hole pairs. The larger energy of these confined localized charge carriers leads to an increased excitonic transition energy [20]. Figure 2-5. Illustration of evolu tion of energy bandgap for bulk cr ystals and nanocrystals [18]. The surface plays an important role in the physical and chemistry properties of nanocrystals, including solubil ity, reactivity, melting point, electr onic structure, surface area, optical emission and chemical stability. For nano crystals in the range of nanometer, a large portion of the atoms are located on or near the surface of nanocrystals. For a 50 CdS nanocrystal, ~15% of the atoms sit on the surfaces [21]. Several methods have been reported to modify the surface to achieve be tter surface physics and chemistry properties. It is generally believed that surface modification passivates uns aturated bonds leading to a reduction of non-

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29 radiative quenching sites, i.e. enhancement of the luminescence efficiency [22-25]. Surface passivation with organic capping ag ent and inorganic shell have been used to stabilize and reduce the localized quantum states from defects that lie in the forbidde n band gap [18]. Various organic/inorganic materials have been utilized as capping layer on the nanocrystals surface through covalent or ionic interactions [26-29] This capping shell not only affects the surface electronic state passivation, but al so may improve the stability and solubility of nanocrystals in desired solvents by preventing aggregation and precipitation [30, 31]. Besides modifying the electronic states of nano crystal surfaces, the cappi ng layer may also facili tate the connection of nanocrystals to the surrounding media [16]. Th ese core/shell nanostructures may provide economic advantages, such as precious materials has deposited on inexpensive core materials via colloidal solution method to achieve reduced cost. The capping layer is also often used to alter the kinetics of nucleatio n and growth, so as to manipulate the anisotropic growth leading to various shape and morphology which will be disc ussed in the following section [32]. A variety of possible nanocrystals core/shell schemes are ill ustrated in Figure 2-6. In each case, the surface of nanocrystals may be modified to meet the demands of an application and to improve the physical and chemical properties. Figure 2-6. Illustration of core /shell structure of nanocrystal s. (a) Surface modified core materials with organic molecules. (b) Complete shell coverage of core materials. (c) Core material with a thin shell layer. (d ) Encapsulation of shell materials into core materials. (e) Hollow core/shell structure. (f) Core/multi-shell structure [33].

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30 2.1.3 Growth Mechanism and Shape Control Several synthesis methods have been deve loped to prepare nanocrystals of various materials as described in the pr evious section. It is still a challenging ta sk to elucidate the mechanism for nano-size crystals due to their hi gh surface-to-volume ratios. In this section, the nucleation and growth mechanisms leading to mono-dispersed distributions and anisotropic shapes are discussed. To achieve a mono-disperse d size distribution, it is necessary to induce a single nucleation event and prevent additional nucleation during the synthesis process. In a burst nucleation process, many nuclei are generated almost simultaneously without additional nucleation afterwards which re sults in mono-dispersed crysta l size distribution. These nuclei appear in a homogeneous solution without other seeds for heteroge neous nucleation. In contrast, if additional nucleation events happen after the first burst nucleation event, size control is not easy to control. Figure 2-7 shows a typical nucleatio n and growth diagram. In stage the concentration of monomer constantly increases with time. However, due to the high energy barrier of spontaneous homogeneous nucleation, nuclea tion does not occur. When the monomer concentration reaches a supersaturation level as shown in stage nucleation occurs. Growth of nuclei into nanocrystals reduces the monomer concentration st opping nucleation. The crystals keep growing until the monomer concentr ation is low, as shown in stage [34].

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31 Figure 2-7. Illustration of nuclea tion and growth diagram. S=Sc: critical saturation of monomer concentration to induce nucleation. S=1: equilibrium monomer concentration below which growth stops [34]. Homogeneous nucleation can be understood based on thermodynamics. The Gibbs free energy for formation of spherical crystals with ra dius R from the solution with supersaturation S is given in Equation 2-1, where is the surface free energy per unit area and is always positive. GV is the free energy change between the monome rs in the solution and a unit volume of bulk crystals and is negative once th e monomer is supersaturated. GvRRG 3 23 4 4 (2-1) The value of R at which G reaches a maximum positive value is called the critical radius RC which indicates the minimum radius of a nucleus that can spontaneously grow larger in the supersaturated solution. The value of RC can be obtained by setting d G/dR=0 and the result is shown in Equation 2-2, where Vm is the molar volume of bulk crystal. Based on Equation 2-1, nuclei with radius greater than RC will not dissolve in the so lution, and the degree of supersaturation should be high to overcome th e homogeneous nucleation barrier. During the growth process, the saturation ratio continuously decr eases and the correspondi ng critical nuclei size continues to increase. In this stage, any pa rticles smaller than this new critical size will disappear through either the disso lution in the solution or the Ostwald ripening process [35].

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32 SRT Vm G RV Cln 22 (2-2) The growth of nanocrystals inevitably involves the process of precipitation of a solid phase solution. As discussed above, the process and parameters can be contro lled during the nucleation stage. Further growth of nanocrystals of desire d size and shape can be manipulated by a number of parameters. Nanocrystals of metal oxide, metal and semiconductor materials of various shapes have been successfully prepared over the last decade, including rods, wires, tetrapods, dots and other unique shapes [32, 36-45]. In this fo llowing section, we will discuss some major mechanisms, monomer concentration and temperat ure, and surface energy modulation, which are commonly used to control the formati on of different nanocrystal shapes. The formation/nucleation of na nocrystals relies on a supersaturated solution which can be achieved either by directly dissolving the solute at higher temperature a nd then cooling to low temperatures, or by adding additional reactants and surfactants to produ ce a supersaturated solution during the reaction. This latter mechanism has been used in tuning the reaction variables, i.e. monomer concentration and reaction temperatur e. The nucleation stage plays a prime role in determining the size and shape of the resulting nanocrystals. For a given monomer concentration and thermal energy in nucleation stage, a critical nanocrystal size can be obtained. The relative chemical potential is highly size -dependent, and very sensitive to the configuration of the nuclei. In general, nanocrystals grow along the crystal di rection with the lowest energy at equilibrium, leading to anisotropic shapes. Figure 2-8 shows an example of anisotropic growth of CdSe as a function of monomer concentrati on. Anisotropic growth of nanoc rystals along specific direction is preferred under the conditi on of high monomer concentra tion with sufficiently high temperatures. In contrast, an isotropic shape is obtained with a low monomer concentration [46].

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33 Figure 2-8. Illustration of the anisotropic growth of CdSe nano crystals by varying the monomer concentration [46]. In addition to controlling monomer concentrat ion and thermal energy (temperature) during nucleation, the anisotropic shape of nanocryst als can be manipulated by modulating the surface energy through introduction of surfactants that absorb onto the surfaces /facets of the growing crystals [31, 40, 47]. At equilibrium, the Gibbs -Curie-Wulff theorem states that the shape of a crystal is determined by those facets which mini mize the total surface energy [48]. EI-Sayed, et al have demonstrated control of the shape of transition metal nanocrystals by surfactants which bond to the surfaces [49, 50]. Figure 2-9 shows an isotropic growth along the [001] direction of CdSe nanocrystal [51]. The surface growth ra te was controlled with selective capping by surfactants on the growing crystal facets. The capping surfactants hinder the growth of the nanocrystal on some facets, lead ing to the formation of nanorods A similar case can be found in Co and CuS nanocrystals where preferential bonding of amine on the {001} facets leads to growth in the [100] and [110] directions and formation of di sk-shaped nanocrystals, as shown in Figure 2-10 [52-55].

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34 Figure 2-9. Anisotropic growth along [001] as a func tion of capping surfactants concentration leading to various aspect-ratios of CdSe nanorods. (a) short nanorods (aspect ratio:5.5), (b) medium nanorods (aspect ratio:11), and (c) l ong nanorods (aspect ratio:20) [51]. Figure 2-10. Anisotropic disk-sha ped nanocrystals of (a) Co and (b) CuS grew along [100] and [110] directions due to selective bonding of the capping surfactant on [001] facets [52-55]. Besides semiconductor nanocrystals, several rare-earth oxide nanocrystals have been synthesized with an anisotropic morphology from formation of metal-surfactant complexes. These complexes organized into layered nanostruc ture as illustrated in Figure 2-11(a). Thus, uniform plate-shaped nanocrystals, such as Eu2O3, ZnO and Nb2O3, were prepared with thickness on the order of a unit cell as shown in Figure 2-11(b)-(d), respectively [56]. In addition, Seo, et al. reported the anisotropic shapes of Euand Tb-doped colloidal Gd2O3 nanocrystals grown non-hydrolytically, using olei c acid, oleylamine, benzyl ethe r and trioctylphosphine oxide. The nanocrystals shape was a func tion of the gadolinium precursor, surfactant and heating rate. Figure 2-12 shows TEM images of the shape of the Gd2O3:Eu3+ nanocrystals versus the

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35 gadolinium precursors. The nanocry stal shape was suggested to result from the dependence of the reaction rate upon crystal or ientation, leading to a variety of shapes [57, 58]. Figure 2-11. (a) Schematic illustration of the self-assembly of rare-earth oxide nanoplates. Examples of the self assembled nan oplates of and the examples on (b) Eu2O3 (c) ZnO, and (d) Nb2O3 [56]. Figure 2-12. TEM images illustrating the dependence of the shape of Gd2O3:Eu3+ nanocrystals on the Gd precursor: (a) Gdacetoacetonate, (b) Gd-acetate and (c) Gd-chloride [57, 58]. 2.2 Photoluminescence Enhancement of Nanocrystals 2.2.1 Surface Passivation Effect Nanostructured materials are building blocks with dimensions typically in the 1 to 100 nm range, and possess properties that are size-depe ndent and often uniquely different from those exhibited by their bulk counterpa rts. Their novel and tunable prope rty are of great interest for

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36 both fundamental understanding of materials and technological a pplications. The effects of surfaces on luminescence from nanoparticles have been frequently reporte d [15, 59]. To achieve stable photoemission with high quantum efficiency, passivation of the surf ace is crucial and is generally achieved with an organic or inorgani c capping layer on the nanoparticles, as shown in Figure 2-13. For inorganic passivat ion layer, a higher band gap material is chosen to confine the electrons and holes to the core, leading to l onger decay lifetimes and higher quantum yields. Figure 2-13. Schematic diagram of inorgani c and organic capping sh ell on a luminescent nanocrystal core [15]. Kmpe, et al have reported an increase of more than 60% in the quantum yield of Tb3+ luminescence in CePO4:Tb nanoparticles due to surface modificat ions [60]. Yang, et al reported photon-stimulated oxidation of the ZnS shell to ZnSO4 over a CdS:Mn nanocrystal led to a 40% increase in brightness [14]. The Mn doped CdS core nanocrystals were formed by a reverse micelle method using Cd2+, Mn2+ and S2rapidly stirred for 10-15 min followed by dropwise addition of Zn2+ solution to form a ZnS shell. Yang, et al. reported significantly enhanced PL from CdS:Mn quantum dots capped with either an inorganic ZnS shell, or from n-dodecanethiol organic molecules, as shown in Figure 2-14. Th e passivation by an organic shell is frequently inferior to that achieved by inorganic shells, and also deteriorates with time. Therefore, surface

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37 defect emission from CdS quantum dots is obser ved and the overall quantum yield of organicpassivated nanocrystals is low (s ee Figure 2-14). In contrast, passi vation of inorganic materials with a large bandgap on the su rface reduces the surface quenc hing sites leading to the enhancement of luminescence [15]. Figure 2-14. The effects of organic and inorga nic shells on the PL spectra and QYs from CdS:Mn luminescent nanocrystal cores [15]. While the effects of surface shell passivati on on semiconductor nanocrystals are significant, enhanced PL has also been reported from nanoc ores doped with rare-earth ions and capped with an undoped shell of another material, e.g. YVO4:Eu3+/YPO4, YVO4:Eu3+/YBO3 and CeF3:Tb3+/LaF3 [23, 25, 61]. Similarly, enhanced PL has been reported for an undoped shell and a doped core, where the shell and the co re materials are the same, e.g. LaPO4:Eu3+/LaPO4, LaF3:Eu3+/LaF3 and Y2O3:Eu3+/Y2O3 [62-64]. The substantial improvement in PL intens ity by surface modification results from the dominance of surface atoms in na noscale particles. Due to their high surface to volume ratio, the density of surface dangling bonds is high on the part icles, and they act as sites for nonradiative

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38 relaxation, reducing the luminescent lifetim e and quantum efficiency. For a spherical nanoparticle composed of 104 atoms (e.g. 4 nm in diameter), the percentage of atoms lying in the outermost atomic layer is about 20%. However, it is more realistic to c onsider the surface of a nanoparticle to be the volume of a thin shell, instead of just the outermost atom layer. In a 4 nm diameter nanoparticle, half of th e atoms will be contained in a 0. 4 nm thick surface shell, making them sensitive to the environment and surface chemical reactions/interactions. Moreover, the fraction of atoms in a surface shell increases inversely proportional to the radius of the nanoparticle, i.e. as the diameter decreases below 4 nm, the majority of atoms are in the 0.4 nm surface shell. 2.2.2 Sensitization for Energy Transfer Effect Energy transfer between two ions requires a certain interaction. In general, electrons on the luminescent centers (activators) were brought to an excited state, followed by return to the ground state through radiative or non-radiative relaxation. During this process, another possibility route to return to the ground state is via the energy transfer from activators to another center, viz. sensitizers. Figure 2-15 illustra tes the process of ener gy transfer through the sensitization effect. The activator is denoted as A, while sens itizers are denoted as S (other literature use D and A for donor and acceptor). If the emission from activators is due to the energy transfer from other ionic species, then they are called sensitizer activators. However, these ionic species may function as quenchers, rath er sensitizers, and diminish the emission from activators. If the sensitizers and activators wave functions overlap, energy transfer can occur. As shown in the schematic, this interaction can be an exchange interaction or an electric/magnetic multipolar interaction.

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39 Figure 2-15. Schematic illustration of energy transfer by the exchange interaction [65]. Resonant energy transfer between the sensitiz er (donor) to the activator (acceptor) can occur when the two ions are within a critical in teraction distance. The close proximity of the S and A centers results in a signifi cant electrostatic interaction, t ypically in the form of electricdipole transition. For a resonant energy exchange, the energy di fferences between the ground and excited sates of the sensitizer a nd activator must be equal. The energy transfer rate equals the radiative rate, if the sensitizer and activat or are separated by a critical distance (RC). For R>RC, radiative emission from the sensitizer dominates, while for R
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40 Tb2S3 nanoparticles. Bang, et al. [68] and Ntwaeabor wa, et al. [69] reported enhanced PL from SiO2:Eu3+ and SiO2:Ce3+, respectively, due to energy transf er from ZnO nanoparticles. Nogami, et al. reported increased PL from Eu3+ in glass by incorporating SnO2 nanocrystal [70]. The sensitization with nanoparticles in the host matrix suggests that the energy transfer processes could occur through similar mechanisms as energy transfer between ions in a host matrix [71]. Nanocrystals may behave like phosphor sensitizers, harvesting the excitation energy from absorption of photons and subsequently transferring it to activators resulti ng in increased PL. For the ex-situ embedding of nanocryst als into the host matrix as discussed above, nanocrystals should be homogeneously distributed into the host matrix without segregation or aggregation in order to effectively transfer ener gy to activator ions. So lid state reaction methods have been used to sensitizing Bi3+ ions with Eu3+, e.g. YVO4:Bi3+, Eu3+, Y2O3:Bi3+, Eu3+, CaWO4:Bi3+, Eu3+, and Y3Al5O12:Bi3+, Eu3+ [72-80]. Wang, et al. also reported that co-doping with Ba2+ ions led to enhanced PL from YVO4:Ln3+ (Ln = Eu, Dy, Sm, and Ce), due to energy transfer from VO4 3to Ln3+ as well as the excited state O2to Ln3+ generated by co-doped Ba2+ ions [81]. Shim, et al. found enhanced PL of laser-ablated GdVO4:Eu3+ thin film by co-doping Li+ ions [82]. 2.3 Inorganic Luminescent Materials Luminescent materials have been widely stud ied and applied in a number of applications including solid state lig hting, bio-labeling, display, photovol taics and scintillators. The small amount of intentionally doped impurity in the host matrix is called a luminescent center or activator. In the following sect ion, rare-earth doped lumines cent phosphor materials will be the main focus and the potential utilization of these materials will be discussed. 2.3.1 Fundamental of Luminescent Materials For extrinsic inorganic luminescent materials, luminescence results from intentionally incorporated activators, usually rare-earth or transition metal ions. The radiative mechanism for

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41 extrinsic luminescent materials is primarily an electron-hole recombination. When external energy is absorbed by the electr ons in the phosphor from a prim ary electron, photon or a charge carrier accelerated by an electric field, the phosp hor electrons are excited to the conduction band. Subsequently the excited electron returns to a lower energy or gr ound state and the energy difference is released as a photon or as phonons in a radiative or non-radiative relaxation, respectively. In a typical photoluminescence process shown in Figure 2-16, an electron in luminescent phosphor materials changes from the ground state to an excited state due to absorption of a photon. The excited electron relaxes to a lower en ergy excited vibrational state via a nonradiative process, followed by radiative relaxation to th e initial state by emitting a photon that is red shifted relative to the excitation photon energy/wavelength. Th e time for this process, called the luminescent relaxation lifetime, is usually in a range of 10-9 ~ 10-3 sec. While phosphorescence occurs through a complex intersystem crossing from singlet state to triplet excited states as shown in Figure 2-16(b) and the lifetime is usually in a range of 10-9 ~ 10 sec [83]. Figure 2-16. Schematic diagram of luminescence from (a) fluorescence, denoted as (f) and (b) phosphorescence, denoted as (p) [83].

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42 Because photoexcited electrons rapidly therma lize and are captured by the states within ~kT of the lowest energy levels, the excited elec trons tend to non-radiativel y relax to the lowest excited energy state. Therefore, an increase in wavelength (decrease in energy) is observed between absorbed and emitted photons. This energy shift is called the Stokes shift as shown in Figure 2-17(a). Luminescence generated from in tentionally incorporated impurities is called extrinsic luminescence. These impurities create local quantum stat es that lie within the bandgap, that result in radiative rela xation by an electron-hole recomb ination from conduction band to acceptor state, donor state to valence band or donor state to acceptor state, as shown in Figure 217(b). Instead of the photon energy being equal to the bandgap, a Stokes shift is observed due to the localized quantum state lying in the forbidden bandgap. These transitions between the excited and lower energy states are classified as allowed or forbidden based on the Laporte and parity selection rules [84]. The electric dipole transition can occur be tween energy levels with different orbital angular momentum, l= and parity. Even though a transition may be forbidden, incorporation of activator ions in crystals may perturb the symmetry and magnitude of the crystal electric field, resulting in obser vation of allowed transition. Al lowed transitions include s-p and f-d, while d-d and f-f are forbidden. The fd transition processes are allowed so optical absorption is strong and obser ved as broad band excitation around 300 nm. The energy of f-d transition is dependent on the local crystal fiel d, while the energy of forbidden f-f transitions from rare earth activators are cons tant for different hosts because of isolation of the 4f states. Luminescent 5d-4f transitions are found for some rare-earth activators, e.g. Ce3+, Pr3+ and Eu2+, and are sensitive to crystal field splitting in host crystals. Forbidden d-d transitions are often observed from activators such as Mn2+, Fe2+ and Cr2+, but luminescent lifetimes on the order of millisecond are observed [14]. In d-d transitions, the d-orbitals extend furthest from the atom

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43 nucleus and are not shielded by a ny other electron orbitals. Ther efore, the energy levels of dorbitals are strongly influenced by the host environment and split by the surrounding crystals field, leading to broad (bandtype) emission peaks. In addition, some trivalent rare-earth ions show f-f transition such as Tm3+, Er3+, Tb3+ and Eu3+. Since the 4f energy levels over rare earth activators are shielded by the outer 5s and 5p electron orbitals, their transition energies are largely unaffected by the host materials and the em ission spectra exhibit sharp atomic-like peaks. Figure 2-17. (a) Stokes shift is the energy differen ce between absorption and photoluminescence emission. (b) Radiative recombination pro cesses in extrinsic luminescence [14]. 2.3.2 Application of Luminescent Materials Luminescence can occur as a result of various excitation source, which is usually denoted as photo-, electro-, cathode-, chemo-, thermo-, s ono-, or tribo-luminescence as shown in Table 21. Luminescent materials can be found in a broad range of everyday applic ations such as cathode ray tubes (CRTs), projection televi sions, fluorescent lamps, X-ray de tectors, solid state lighting, sensors, and displays. Consequently, research and development on luminescent materials has resulted in synthesis and test ing of thousands of phosphors. Ho wever, only about 50 materials exhibit properties that are suit able for appropriate technologi cal applications in terms of

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44 efficiency, emission color, decay time, physical stability, availability of raw materials, environmental aspects, cost, reproducibility and ease of materials preparation [83, 85, 86]. Table 2-1. Type of luminescence gene rated by different excitation sources. Luminescence type Excita tion source Application Cathodoluminescence Electrons TV sets, monitors Photoluminescence UV (photons) Fluorescent lamps, plasma displays X-ray luminescence X-rays X-ray amplifier Electroluminescence Electric field LEDs, EL displays Sonoluminescence Ultrasound Chemoluminescence Chemical reacti on energy Analytical chemistry Triboluminescence Mechanical energy Todays efficient, low-cost fluorescent lighti ng is mainly based on i ndirect light emission from mercury plasma discharges. The main em ission from a mercury discharge can be tuned from 185 to 254 to 365 nm when the mercury pressure is increased. In fluorescent lamps, a combination of phosphors emitting at the appropriat e colors are coated on the inside of the discharge tube, as shown in Figure 2-18. Select ion of the appropriate luminescent materials enables special application of fluorescent lamps su ch as water disinfection, security marking, sun-tanning, photocopying, display backlight an d advertising billboards. An increasingly important application of luminescent material is for down conversion phosphor to create white light emitting diodes (W-LEDs). White light can be generated by combining the output from blue, green and red-emitting diodes to yield wh ite light. However the most common method to generate white light is to use a fraction of the light from a blue LED to excite emission from a phosphor(s) that yield green, yellow and/or red light, as shown in Figure 2-19. Highly stable, yellow-emitting yttrium-gadolinium aluminum-g allium garnet doped with cerium, (Y,Gd)3(Al,

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45 Ga)5O12:Ce3+, is used in commercial W-LEDs in traffi c lights, car headlights, outdoor lighting, flashlights and display lighting [87, 88]. Figure 2-18. Schematic diagram of a mercury-discharge fluorescent lamp illustrating the light generation procedure and possible emission colors [89]. Figure 2-19. (a) Schematic diag ram of blue LED with YAG:Ce3+ phosphor coating. (b) Emission spectrum illustrating a combin ation of blue LED and YAG:Ce3+ phosphor coating [90]. 2.4 Inorganic Scintillation Materials Scintillation is an example of radioluminescen ce. Ionization radiation or particles will be absorbed and a fraction of the energy will be conve rted into visible or ultraviolet photons. This conversion process typically takes place on a time scale of nanoseconds to microseconds, producing a fast pulse of phot ons corresponding to each absorp tion event by the scintillator material [91]. The number of light pulses and in tensity are proportional to the type of particle

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46 and the energy deposited in the scintillator, wh ich is further processed by converting the photons signal into an electrical signal. Scintillators can be gaseous, liquid or soli d, organic or inorganic and crystalline or amorphous. Here, solid state inor ganic scintillation materials will be the focus due to their excellent characterist ics for detection of radiation. 2.4.1 Development and History of In organic Scintillation Materials In 1895, Roentgen discovered X-rays by the fluorescence excited from a barium platinum cyanide scintillator phosphor. Howeve r, he realized that this t ype of material was not very efficient absorbing the X-ray and converting it in to visible emission. He demonstrated that cadmium tungstate (CdWO4) was much more effective in converting absorbed X-rays into visible emission recorded on photographic films [ 92]. The discovery of visible light emission from certain materials when expos ed to radiation led to the use of scintillators as radiation detectors. A large number of compounds were synthesized and tested for fluorescence under Xray irradiation. Scintillators have been played a major role in th e development of modern physics. E. Rutherford observed particles by scintillation from a zinc sulfide screen. Until the end of the World War ZnS and CdWO4 were the most popular particle sc intillator detector materials in nuclear physics. Data manipulation techniques fr om related detection methodologies, such as electronic counting methods from gas ionizat ion chambers, were adopted by scintillation detectors. In 1947, sodium iodide doped with thallium (NaI(TI)), a widely used scintillation materials nowadays, was discovered in this period. In 1948, Hofstadter used polyc rystalline NaI:TI to detect particles, and soon single cr ystals were coupled with a PMT to provide spectroscopic information on the energy and type of radiation source. Prior to this discovery, spectroscopic information could only be obtained by inefficient methods such as reflec tion by crystal lattice or with proportional gas ionization detectors. Spectr oscopic information was a huge achievement in

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47 the field of radiation detection and it paved th e way for further scinti llator development. Many compound materials were reported to be good scintillators during the following two decades, such as CsI:Na, CsF, CaF2:Eu and Bi4Ge3O12. In the 1980s, significant progress in scintillator materials was made due to the demanding requirements in medicine diagnosis and hi gh energy physics. Development of cerium-doped lanthanum halides in 1997, including LaCl3:Ce and LaBr3:Ce scintillator single crystals, yielded new benchmarks in high light yield and excellent energy re solution [93-95]. However, detrimental properties such as hygroscopicity, ease of cracking, limited size and high cost limited their application. The search and study for new scintillation materials remains a challenge even today. In addition to development of new scintill ator materials, a great deal of research is focused on new synthesis methods and detector configurations to reduce the cost. Figure 2-20 shows an evolution of the important discoveries fo r inorganic scintillator materials over the last century [96]. Figure 2-20. Timeline for the discovery of impor tant inorganic scintillator materials [96].

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48 2.4.2 Fundamentals and Criteria of Scintillation Materials Scintillator detectors are composed of the scintillator (phosphor) material, optical coupling components and a photodetector, as shown in Figur e 2-21. After an energetic photon or charged particle is absorbed, the ultraviolet, visible or near infrared light generated by the scintillator conversion is coupled to the photode tector and the output electrical signal is further processed, typically by an amplifier and multi-channel pulse height analyzer. The scintillation peak is proportional to the energy of the radiation. Th e most common display mode of the radiation spectrum is the differential pulse height distribu tion, as shown in Figure 2-22(a). This spectrum may include peaks from photoelectric and Compton e ffects [97] in addition to the main radiation peak. The pulse height distribution depends on th e relative cross sections for the absorption and excitation processes. The energy resolution define s the detectors ability to distinguish between radiation events with very close energies. The ab ility of a scintillator to resolve two gamma rays with different energies increases as the energy resolution improves. Energy resolution is defined to be the full width at half-maximum (FWHM) di vided by the height of the peak, as shown in Figure 2-22(b) which is a spectrum of 55Fe measured with a cerium doped yttrium aluminate scintillator. Figure 2-21. Schematic diagram of a typical scintillator detector.

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49 Figure 2-22. (a) Differential pulse height distribution for gamma rays from 137Cs source. (b) The scintillation pulse he ight spectrum from 55Fe measured with a YAP:Ce scintillator at room temperature. The first stage in the scintillation detection proc ess is the conversion of kinetic energy into photon emission. The physical processes leading to scintillation in inorganic solids have several stages including relaxation of the initial elec tronic excitation, therma lization, trapping of electrons and holes, and radiativ e relaxation of the luminescent centers as shown in Figure 2-23 [91, 96]. During the initial absorp tion process, a number of inte ractions can occur between the high energy photons (considered <1 MeV) and the lattice of scint illator materials, e.g. the photoelectric effect and Compton scattering e ffect. When radiation is absorbed by the photoelectric effect, a photoelectro n and an X-ray are released. Compton scattering refers to the inelastic scattering of photons from free or l oosely bound electron which are at rest. Absorption of a high energy photon creates a single high en ergy electron which may be scattered by the electrons of the host, creating many hot electron-hole pairs, which may be thermalized by the intraband transitions so that the electrons a nd holes relax to the bottom of the conduction band and the top of the valence band, respectively. Electrons and hol es migrate through the materials and either become trapped on impurities, self-trapped by crystal lattice, or combine by Coulomb

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50 attraction to form free and impurity bound excitons in a time scale of ~10-12-10-11 sec. Due to capture at trapping levels in the forbidden ga p of the materials, non -radiative relaxation may degrade the scintillation materials performance. For the final stage, the excited luminescent species may return to a lower energy (often th e ground) state either by non-radiative quenching processes or by radiative recombinations to em it photon. The radiative process can be short as 10-9 sec for electron-hole recombination, but is typically slower for emission from free and bound excitons (10-6 sec to minutes). Figure 2-23. Illustration of the m echanisms involved in the scinti llation process in an inorganic material [96]. There are various luminescent species and radiative processes i nvolved in inorganic scintillators. Scintillating materials can be di vided into two major categories, intrinsic and extrinsic scintillators, as shown in Table 2-2. Luminescence from intrinsic or self-activated scintillators results from different types of radi ative transitions, such as recombination of an electron with a hole, free (self-trapped) or impur ity-trapped excitons, charge transfer (within a molecular complex) and core-valence band transitions.

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51 Table 2-2. Selected scintillator material s and their radiative transitions [91, 92]. Category Scintillation materials Emission transition Intrinsic (self-activated) CeF3 Bi4Ge3O12 CsI, NaI, LaF3 CdWO4 CsF, BaF2 5d-4f 6p-6s self-trapped exciton charge transfer core-valence band Extrinsic (activated) Gd2O3:Eu3+, Gd2O2S:Pr3+;Tb3+ CaF2:Eu2+ Lu2SiO5:Ce3+, LaF3:Nd3+ NaI:TI+, CsI:TI+ CsI:Na+ ZnS:Ag, PbI2 4f-4f 5d-4f 5d-4f 6p-6s excitonic donor-acceptor pair For free or impurity-trapped exciton lumine scence, ionization electrons and holes may create excitons with binding energy of 50~100 meV. At lo w temperatures, these excitons remained bound and decay radiatively with a short time constant (<1 ns). However, at room temperature the emission from these excitons ma y be weak due to thermal dissociation. In the case of self-trapped excitons, the hole localizes on one or more atoms resulting in a lattice relaxation. The resulting distortion traps a spat ially diffuse electron to create a self-trapped exciton. A triplet state is formed by unpaired spin associated with the hole and the diffuse electron, and luminescent relaxation is forbidden by the parity selection rule resulting in a long radiative lifetime is in a range of 10-6~10-3 sec. The singlet state can decay much faster with lifetimes on the order of 10-9 sec. Some representative examples include NaI, CsI, BaF2 and LaF3. In other self-activated intrinsic materials, the luminescence may result from material components, such as in traionic transitions on Bi3+ in Bi4Ge3O12 and Ce3+ in CeF3, or from charge transfer transitions on (WO4)2in CaWO4 and CdWO4. Finally, intrinsic luminescence may result from core-valence band transitions in materials such as BaF2 and CsF. A photon is emitted when an electron in the valence band fills an ionization hole in a shallow core level, especially when the energy difference between the valence band a nd core level is less than the bandgap. The

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52 lifetime of this transition is short (~1 ns) but the light yield is low due to the inefficient creation of holes in the shallow core level. For activated extrinsic luminescence, activators (dopant) ions are incorporated into the materials leading to localized quantum states which trap the ionized electrons and holes. The activators may result in luminescence directly or may promote luminescence from defect-bound excitons. The electric-di pole 5d-4f transition on Ce3+ is governed by the spatial overlap of 5d and 4f orbitals. The transition is allowed so the lifetime is typically fast (~ 20-60 ns). In contrast, the lifetime of the Eu2+ 5d-4f transition is slow (~ 1 s) due to the partially spin allowed transition. For lanthanide ions such as Eu3+, Pr3+ and Tb3+, the electric-dipole 4f-4 f transition is forbidden resulting in a long lifetime (~ 0.1-1 ms). To be a good scintillation material, several physical, optical and ch emical properties are required as listed in Table 2-3. There is no perfect scintillator material suitable for all, since different applications require di fferent priorities and demand. Ther efore the selected scintillation materials have an optimum combination of prop erties which match the requirements of a given application. Table 2-3. The desired propertie s of an ideal scintillator. Criteria Scintillation properties Detection efficiency High density and atomic number High count rate capability Short decay time; low afterglow Good spatial resolution High light yield Good energy resolution Clear id entification of energy event Suitable emission spectra Good match to PMT or photodiode Good transmission for light to photodetect or Transparent at emission wavelength Radiation hard Stable performance to ionizing radiation Chemical, thermal and mechanical stability Rugged and nonhygroscopic; ease of crystal configuration Low cost fabrication Low cost raw materials; different growth methods

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53 2.4.3 Synthesis Processes and Applicat ion of Scintillation Materials The vast majority of scintillator detectors use single crystals, because of their higher light yield and lower scattering effects. Czochralski or Bridgman growth [ 98] are the most common methods to grow single crystals as long as th e material forms low volatility melts with low thermal expansion and few solid state structural transformations. Growth of structurally or chemically complex crystals with this process is difficult due to factors such as thermal expansion, thermal conductivity, and tendencies to cleave or slip. Moreover, incorporation of dopant ions uniformly and on the proper substitu tion sites requires well controlled conditions during the process. If the desi red compound has a high vapor pressure, then closed vessels are required. Growth configuration can be vertical and horizontal, and may include rotation or titled ampoules because of viscosity and convection of the melt. The complexity of growing large single crystals of scintillators results in a high cost and limite d supply. Polycrystalline ceramic scintillators, such as (Gd,Y)2O3:Pr, have been densified by sinter ing or by hot isostatic pressing to improve the optical propertie s and radiation absorption and we re used for X-ray imaging in CT system. Sol-gel, solid state reaction, hydrothermal and hot-solution methods have been used to grow ceramic scintillation materials. The te mperatures required to make polycrystalline ceramic scintillator are much lower than those for growing single crystals. In addition, the size of the ceramic scintillators is not limited by thes e synthesis methods. However, light scattering, transparency, and thickness are typical of issues in using ceramic scintillators [91, 92]. Scintillators have a number of applications including medical imaging, security, high energy particle physics and non-destructive flow detection. In medical imaging, positron emission tomography (PET), single photon emissi on-computed tomography (SPECT), X-ray computed tomography (CT) and planar X-ray sc reens all use scintill ator detectors. The

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54 scintillators for CT require low afterglow, while PET demands scintillator s with high density and atomic number, fast decay a nd high light output. SPECT requ ires high light output and good energy resolution [99-101]. Figure 2-24 shows the principal components of a typical CT medical imaging application. An X-ray beam rotates around the patient body, along with a 1D-positionsensitive detector, as shown in Figure 2-24(a) An Image, as shown in Figure 2-24(b), is reconstructed from the cross-sectional data. Ceramic scintillators, such as Gd2O2S:Eu, Pr and (Gd,Y)2O3:Pr, have been utilized in CT syst em because of their homogeneity and good machinability, but co-doping was necessary to reduce the afterglow. Figure 2-24. (a) Schematic diagram of computed tomography (CT) imaging. (b) A CT scan image of a transversal slice of human body at the level of the kidneys [99]. Medical imaging by PET is used to characte rize metabolism, for example by injecting two 511 KeV quanta collinearly, such that the positron is annihilated in some particular tissue. The two quanta are detected by coincidence counting us ing position sensitive dete ctors. In general, a PET system consists of many rings with thousands of scintillation detectors, as shown in Figure 2-25(a). Therefore, a 3D image of processes in the body can be rec onstructed, as illustrated in Figure 2-25(b). The most commonly used scintillator in a PET system is Bi4Ge3O12 (BGO) due to its high density, high atomic number and large probability of the photoelectric effect for 511

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55 KeV quanta. However, the energy resolution and light yield of the scintillator need to be improved. Figure 2-25. (a) Schematic diagram of positr on emission tomography (PET) with the inset illustrating BGO detectors coupled with PMT. (b) A PET scan image of a brain from Alzheimers disease patient [102, 103].

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56 CHAPTER 3 SYNTHESIS AND CHARACTERIZATION OF LUMINESCENT ZERO AND ONE DIMENSIONAL GADOLINIUM OXIDE NANOCRYSTALS 3.1 Introduction Rare-earth doped luminescent nanocrystals are of interest due to their potential applications in displays, solid-state lighting, scintillation detectors, and biologic diagnostics [59, 104, 105]. It is well established that the electrical, optical and chemical properties of these shape-controlled or self-assembled nanostructures depend str ongly on crystals structure, morphology and composition [106-108]. In addition, light emitted from rod-shaped nanocrystals is polarized, which may lead to new applications [18, 109]. Zeroand one-dimensional (1D) nanostructures [32, 38, 110], such as quantum dots, nano-rods, nano-wires, nano-tubes and nano-belts, have been widely studied. Specific or ganic surfactants, which can modul ate the growth kinetics, have been employed to grow anisotropic luminescent nanostructures using non-hydrolytic procedures with and without templates [42, 111, 112]. Gadolinium oxide (Gd2O3) exhibits good luminescent pr operties when doped with rareearth ions. The characteristics of the lumine scent and paramagnetic multifunctional properties lead to a number of possible appl ications such as neutron captur e therapy targets and bimolecular detection [113-115]. Moreover, this oxide phosphor is a potential scintillation detector material due to its short luminescent decay lifetime, la rge atomic number and density, low afterglow, large light output and low hygrosc opicity [92, 96]. A variety of s ynthesis methods, such as solid state reaction, hot solution, hydrothermal and combustion met hodologies [1, 57, 116, 117], have been used to grow doped rare-earth nanostructures. In this study, Gd2O3:Eu3+ scintillation nanocrystals have been grown using a variety of synthesis methods, such as sol-gel solution precipitation and hot solution injection growth. The size and morphology of nanocrystals varied

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57 from ~10 nm nanoflowers, ~200 nm nanoparticle s to ~800 nm nanorods. The structural and photoluminescent properties will be illustrated an d discussed. 3.2 Experimental 3.2.1 Non-Hydrolytic Hot Soluti on Phase Synthesis Method In our one-pot hot-solution phase synthesis for the growth of Gd2O3:Eu3+ nanocrystals, Gdacetylacetonate (2 mmol) and Eu-acetate (0.4 mmo l) were mixed and dehydrated at 120 for 1 h, followed by additions of oleic acid (6 mmol), oleylamine (3 mmol), benzyl ether (10 mmol) and hexadecanediol (2 mmol). The precursor solution was magnetically stirred under N2 inert gas flow in a three-neck flask surrounded by a heating mantle as shown in Figure 3-1. The precursor solution was further heated to 120 for 30 min, resulting in a tr ansparent yellowish solution. The temperature increased to 290~320 with a heating rate of 5~25 /min and the condition maintained for 5 to 150 min. The color of the precursor solution changed to deep brown. The solution was sampled at selected intervals to determine the time dependence of growth. The samples were cooled and added to ethanol fo r washing. The procedure of precipitation, centrifugation and washing was repe ated several times to remove residual organics. Purified Gd2O3:Eu3+ nanocrystals were easily disp ersed in non-polar solvent su ch as hexane, chloroform or toluene. Structural and optical properties we re characterized either as solid state powder or colloidal solution.

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58 Figure 3-1. Flow chart of one-pot non-hydrolytic hot solu tion synthesis procedure. 3.2.2 Aqueous Sol-Gel Precipi tation Synthesis Method All chemicals were analytical grade and used as received wi thout further treatment. In a typical sol-gel precipitat ion synthesis of Gd(OH)3:Eu3+ nano-rods, 1.8052 g of gadolinium nitrate (99.9%) was mixed with 0.3568 g of europium n itrates (99.9%) and 100 mL of distilled (DI) water. Nitric acid (6 mL, >62%) was added to th e precursors and the mixture stirred to form a clear solution. The precursor mixture was heated to 90 and 24 g of polyethylene glycol-8000 (PEG-8000) was added into the solution with vigo rously stirring. Then 100 mL of NaOH (4M) solution was added to the precursor mixture with continuous stirring to generate a colloidal dispersion, signified by a cha nge from clear to cloudy. Afte r various reaction times, the precipitate was collected by washing with DI water and centrifuging several times. The washed precipitate was oven-dried at 80 for 2 h, followed by an 800 calcination in air for 2 h. Micro-rods (with larger aspect ratio) were synthesized with a similar procedure except PEG was not added. 3.2.3 Characterization The crystal structure of as-prepared and cal cined samples were characterized by X-ray diffraction (XRD) (Philips APD 3720) with Cu K radiation source (=0.15418 nm). The XRD pattern was collected from dried powder sample s in the step scan (0.02) mode over a 2 scan

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59 range of 20-70. The rods morphology and size were determined with a JEOL 2010F high resolution transmission electron microscope (HR-TEM) operated at an accelerating voltage of 200 kV, and with a JEOL 6335F field emission scanning electron microscope (FE-SEM) operated at 15 kV. The TEM samples were prepar ed by drop-casting rods dispersed in ethanol onto a carbon-coated holey copper grid, fo llowed by drying at room temperature. Photoluminescence (PL) and photoluminescence excitation (PLE) spectra and luminescent lifetime were measured at room temper ature using a JASCO FP-6500/6600 fluorescence spectrometer with a 150 W xenon lamp. Radiolumin escence (RL) spectra were obtained at room temperature using a Moxtek 40 keV, 100 A Magnum x-ray source and an Ocean Optics USB2000 miniature fiber optic spectrometer. For each measurement, a crucible with 57 mm2 area was filled with nanopowder. 3.3 Results and Discussion 3.3.1 Hot Solution Synthesis of Zero Dimensional Gd2O3:Eu3+ Nanocrystals 3.3.1.1 Shape control of Gd2O3:Eu3+ nanocrystals Non-hydrolytic solution synthetic methods are used to prepared nanocrystals or quantum dots with sizes that can be prec isely manipulated in the 1~10 nm range. The control of shape or morphology of these colloidal nanocrystals is possible and has been reported by different synthesis approaches. In the study, Gd2O3:Eu3+ nanocrystals exhibited strong photoluminescence with the peak located at 609 nm a ssociated with th e electric-dipole 5D0-7F2 transition. In addition, these nanocrystals obtained by non-hydrolytic thermal reaction at elevated temperature showed good crystallinity, monodispersed size and hi gh luminescent efficiency. The shape or morphology of nanocrystals was tuned by varying th e experimental conditions such as reaction temperature, reaction time and precursor reagents.

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60 Table 3-1 indicates the experimental parameters used to synthesize of Gd2O3:Eu3+ nanocrystals. Based on these data, nanocrystal s shape can be manipulated by adjusting the reagent concentration, reaction pr ocedure and temperature. In this study, two distinct shapes of nanocrystals, flowerand sphere-like nanocrystal s were prepared by this one-pot hot solution synthesis method as shown in the high-resolution TEM micrographs of Figure 3-2. Figure 3-2(a) shows flower-like nanocrystals with the dimension of ~20 nm. Each nanocrystal is composed of several petals that self-assembled into an indi vidual bigger nanocrystal. In contrast, sphere-like nanocrystals are shown in Figure 32(b) with the dimension of 10 nm. In the process, nucleation is a critical step and it occurs when the monomers concentration reaches supersaturation levels as discussed in chapter 2. The mono mer is continuously incorporated onto the nuclei leading to the growth of the crystals. During th e nucleation and growth stages, the final shape and size of the nanocrystals can be controlled. For sphere-like na nocrystals, there is no preheat procedure during the synthesis. The temperature increased directly from room temperature to the reaction temperature of 320 resulting in a large density of nuclei bursted simultaneously. These monodispersed nuclei continuously grow to form final sphere-like nanocrystals. However, the preheat procedure of the precursor solution at 200 for 30 min dwell time leads to the first stage nuclei at the beginning. The continuous growth reaction to higher temperature produces secondary nucleation process along w ith the continued growth of nuclei of the first stage. Due to the Ostwald ripening effect, the final resulting size is bigger and the shape becomes irregular leading to the flowe r-like nanocrystals.

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61 Table 3-1. Experimental paramete rs of flowerand sphere-like Gd2O3:Eu3+ nanocrystals. Nanocrystals shape Preheat condition Synthesis temperature Quantum yield Flower 200 30 min 290 24% Sphere No 320 87% Figure 3-2. High-resolution TEM images of as-prepared Gd2O3:Eu3+ nanocrystals synthesized with different experimental parameters le ading to nanocrystals with shapes of (a) flowerand (b) sphere-like nanocrystals. 3.3.1.2 Structural and luminescent properties of Gd2O3:Eu3+ nanocrystals Figure 3-3 shows the XRD spectra of as-prepared sphere-like Gd2O3:Eu3+ nanocrystals. XRD patterns show the characteristic peaks of the cubic Gd2O3 crystal phase (JCPDS file #: 431014) with broadened full width at half maximum (FWHM) due to the nanocrystalline size. The result shows that Gd2O3:Eu3+ nanocrystals synthesized via the hot-solution synthesis method are well-crystallized without post-growth heat treatment.

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62 Figure 3-3. XRD pattern of sphere-like Gd2O3:Eu3+ nanocrystals. Figure 3-4(a) shows the Photoluminescence ex citation (PLE) spectrum observed at 609 nm from as-prepared sphere-like Gd2O3:Eu3+ nanocrystals at room temp erature. A broad excitation band from 230 to 300 nm is a charge transfer band (CTB) associated with the 2p orbital of O2to the 4f orbital of Eu3+, whose strength is related to the O-Eu covalency boding and the coordination environment around Eu3+. Another CTB at ~280 nm is related to Gd3+ Eu3+ energy transfer. The weak sharp peaks from 380 to 525 nm are associated with direct excitation of f-f shell transition of Eu3+ which are less dependent of the lattice structure or crystal size of the matrix. Figure 3-4(b) shows the photoluminescence (PL) emission spectrum excited at 260 nm from as-prepared sphere-like Gd2O3:Eu3+ nanocrystals at room temp erature. The photos in the inset show a strong red emission from solid state powder and colloidal solution from Gd2O3:Eu3+ nanocrystals. The red emission from Eu3+ is dominated by the 5D0-7F2 transitions at 609 nm with other minor peaks from 5D0-7Fj (j=0, 1, 3, 4) characteristic peaks. The Eu3+ f-f transitions relative peak intensity are sensitive to the crystal environment, therefore the luminescence can be used as an index in determining the local crystal structur e. The higher intense peak at 609 nm than at 620

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63 nm suggests that Eu3+ spectrum originated from ions on the C2 site in the cubic structure consistent with the XRD data shown in Figure 3-3. Figure 3-4. (a) Photolumines cence excitation (PLE) spectrum of as-prepared sphere-like Gd2O3:Eu3+ nanocrystals for the emission peak at 609 nm. (a) Photoluminescence (PL) spectrum sphere-like Gd2O3:Eu3+ nanocrystals excited by 260 nm. The insets show red emission from solid state powder and colloidal solution samples under UV excitation. Figure 3-5 shows the radioluminescence (R L) spectrum of as-prepared sphere-like Gd2O3:Eu3+ nanocrystals irradiated by 40 keV X-ray. The spectrum is dominated by a peak at 612 nm from the 5D0 7F2 transition similar to the photol uminescence spectrum shown in Figure 3-4(b). Several radioluminescence materials have been already deve loped, e.g. NaI, but most of them are not refractory and cannot be used at elevated temperatures. In this study, a good radiation-induced photoluminescence was observed using Gd2O3:Eu3+ nanocrystals synthesized by this facile one-pot non-hydrolytic synthesis me thod, which supports the pot ential utilization in scintillator detection application.

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64 Figure 3-5. Radioluminescence spectr um of as-prepared sphere-like Gd2O3:Eu3+ nanocrystals. 3.3.2 Sol-Gel Precipitation Synthesis of One Dimensional Gd2O3:Eu3+ Rods 3.3.2.1 Synthesis of Gd2O3:Eu3+ micro-rods In this study, rod-like Gd2O3:Eu3+ nanophosphors and microphosphors were prepared by a sol-gel precipitation method. Rod-like Gd(OH)3:Eu3+ were synthesized at temperatures near 85 then converted to cubic Gd2O3:Eu3+ by calcining at 800 for two hours. The dimension of rods can be controlled by the PEG-8000 which wa s used as a capping agent. The crystal and structures were correlated with the morphology and size of these 1D phosphors as well as the luminescent properties excited by ultra-vi olet (UV) and x-ray irradiation. To investigate the time-dependence morphology evolution of this 1D phosphor, a growthtime analysis was performed with SEM. The morphology and size of the Gd(OH)3:Eu3+ microrod structure as a function of reaction time from 3 min to 2 h are shown by the SEM micrographs in Figure 3-6. The images clearly show the morphology evolved from spherical nanoparticles (6090 nm diameter) after 3 min of reaction time (F igure 3-6(a)) to a mixture of spherical and

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65 rod-like particles after 30 min (F igure 3-6(b)) to predominantly 1D micro-rods with aspect ratio ~10 after 2 h (Figure 3-6(c)). Figure 3-6. SEM microgra phs of sol-gel Gd(OH)3:Eu3+ at reaction times of (a) 3 min showing nano-size spheres, (b) 30 min showing a mixture of sphereand rod-like particles, and (c) micro-rods after 2 h. 3.3.2.2 Capping agent effect Figure 3-7(a) and (b) show SEM micrographs of Gd(OH)3 nano-structures with PEG added and after reaction times of 15 min and 2 h, resp ectively. The morphology evolution from spheres to rods with increasing reaction time was similar to that of micro-rods. However, after reacting for 2 h the rod length was only ~150 nm with a diam eter of ~70 nm similar to micro-rods, i.e. the aspect ratio is only 2. PEG is an amphipathic non-ionic surface active agent from oxyethylene polymerization and therefore a PEG-OH bond shoul d form in aqueous solution and chelate bismuth ions, which would lead to highly mob ile molecules with a large exclusion volume. These PEG-adsorbed nuclei slow the subsequent growth, which leads to nano-size structures with low aspect ratios [118-120].

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66 Figure 3-7. SEM micr ographs of Gd(OH)3:Eu3+ nano-structures synthesized by sol-gel with PEG capping agent for reaction times of (a) 15 min and (b) 2 h. 3.3.2.3 Structural analysis Figure 3-8(a) shows XRD pattern of as-prepared micro-rods grown for 2 h. The diffraction patterns match with hexagonal crystalline Gd(OH)3 (JCPDS card #83-2037; Figure 3-8(b)), even for 3 min reaction time. Af ter being calcined at 800 for 2 hr in air, the diffraction patterns observed for the samples of 2 h (Figure 3-8( c)) are well matched with cubic crystalline Gd2O3:Eu3+ (JCPDS card #43-1014; Figure 3-8(d)). XRD patterns fr om 2 h as-prepared and calcined nano-rod samples show result similar to those from micro-rods. The phase of asprepared nano-rods samples is hexagonal Gd(OH)3 which is converted by calcining to cubic Gd2O3.

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67 Figure 3-8. XRD patterns of as-prepared and calci ned micro-rods samples grown for 2 h. (a) Asprepared samples; (b) Diffraction pe aks from JCPDS card #83-2037 for hexagonal Gd(OH)3; (c) Calcined samples; (d) Diffrac tion peaks from JCPDS card #43-1014 for cubic Gd2O3. Figure 3-9(a) shows a TEM micrograph of as-p repared micro-rods grown for 2 h and the dimensions are consistent with SEM images s hown in Figure 3-6. As shown by the inset of Figure 3-9(a), the selected area electron diffrac tion pattern (SAED) clearly exhibits spotty diffraction rings from hexagonal Gd(OH)3 {100} and {110} planes. The corresponding high resolution TEM (HRTEM) image reco rded along the [110] directi on is shown in Figure 3-9(b). Lattice fringes are clearly visible with a spacing of 0.316 nm which corresponds to the interplanar spacing of the (110) pl ane, indicating the rod axis and growth direction is the [001] direction which lies in the (110) plane.

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68 Figure 3-9. (a) TEM microgr aph of as-prepared Gd(OH)3:Eu3+ micro-rods grown for 2 h. (b) High-resolution TEM image (Scale bar: 5 nm ); the inset is a selected area electron diffraction pattern (SAED). As-prepared and calcined micro-rods grown for 2 h are compared in Figure 3-10(a) and 10(b), respectively, and the tran sformation from hexagonal Gd(OH)3 to cubic Gd2O3 did not change the rod morphology or dimensions. This transformati on to oxides upon calcining has been reported for other lanthanum hydroxide ma terials [121] and the mo rphologies were also maintained after calcining, presumably due to lo w driving forces and/or high kinetic barriers to changes in these anisotropic structures. Figure 3-10. SEM micrographs of 1D micro-rods of (a) as-prepared hexagonal Gd(OH)3:Eu3+ and (b) calcined cubic Gd2O3:Eu3+. 3.3.2.4 Formation mechanism In this work, the growth of Gd(OH)3 with a rod-like morphology was not assisted by a catalyst or directed by a template. The nanoor micro-rods were synthesized at relatively low

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69 temperature (~85 ) in times of 30 min with or without, respec tively, the PEG capping agent. The rod-like morphology presumably results from a nucleation and anisotropic growth process. Colloidal Gd(OH)3 particles were nucleated when NaOH wa s added to the precursor solution. Apparently, the kinetics of growth are rapid along the c-axis and slower along the a and b axes of the hexagonal Gd(OH)3.. Addition of the PEG capping agent apparently slows the growth along the c axis more than it does along other axes, leading to an aspect ratio of 2 for the nano-rods versus 10 for the micro-rods. Based on the Gibbs-Curie-Wulff theorem, the shape of crystalline solids can be explained by the relative growth kinetic s of each facet or phase of the crystals [48]. Therefore, the introduction of suitable capping ag ent into the reaction system can tailor the growth kinetics by modulating th e relative free energies of va rious crystallographic facets. Presumably, inhibition of growth along the c axis results from adsorption of PEG on the hexagonal basal plane which limits the access of precursors to the surf ace reaction sites. A similar mechanism was reported in the growth of ZnO nano-cones and FeCo nanocubes in the presence of PEG [122, 123]. 3.3.2.5 Luminescent and scintillation properties of Gd2O3:Eu3+ nanocrystals The calcined samples showed a bright red emission from the f-f transition of the 5D0 to the 7FJ states for Eu3+ upon excitation by a mercury-discharge ultra-violet (UV) lamp. Photoluminescent excitation (PLE) and photoluminescent emission (PL) spectra are shown in Figure 3-11(a) and 11(b), respectively, for bot h microand nano-rod phosphors. In the PLE spectra in Figure 3-11(a), the broad peak at < 300 nm is from the oxygen to europium charge transfer band (CTB) which is strongly de pendent on the lattice structure of the Eu3+ matrix [1]. The sharp peaks between 360 and 530 nm are from europium self-excitation which are generally less dependent of the lattice structure or crystal size of the matrix, thus they occur at the same

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70 wavelengths for microand na no-rod phosphors. From the PLE sp ectra in Figure 3-11(a), the CTB is much less intense for nano-rod as compared to micro-rod samples and the peak position is 253 and 248 nm for microand nano-rods, respectively. Due to the weak CTB, the Gd2O3:Eu3+ nano-rod samples show lower PL emission (Figur e 3-11(b)) even though th e PL spectra for the 5D0 to 7Fj europium transitions are similar and are dominated by the peaks at 612 and 627nm. The Eu3+ ions precursor concentration is 20 at% with respect to Gd3+ ions in the nanoand micro-rods and the actual incorporation level determined by energy dispersive spectrometry (EDS) is ~18 at% for both calcined nanoand micro-rods. The lower PL emission from nanorods could also result from a larger non-ra diative recombination rate due to a larger concentration of defects based on their larger surface to volume ratio which was reported in other hydrothermal synthesized 1D rods [2]. In order to evaluate the effects on PL intens ity of surface defect st ates on nanoand microrods, photoluminescence decay lifetime data were monitored and analyzed for the 5D07F2 transition at 612 nm as shown in Figure 3-11(c) and 11(d), respectively. The experimental decay profile can be fitted by a single exponen tial function as shown in Equation 3-1. I(t) = I0 exp(-t/ ), (3-1) where I0 is the initial emission intensity at t=0 and is the lifetime of the emission center. The lifetimes for nanoand micro-rods were 0.63 (0.04) and 1.19 (0.06) ms, respectively. The lifetime for the luminescence of Eu3+ is consistent with the reported values for other Eu3+ doped Gd2O3 system [1, 124]. The shorter lifetime for nano -rods can be attributed to the non-radiative relaxation due to quenching centers and defects surrounding europium ions. The shorter lifetimes for nano-rods are consistent with lower PL intensities from nano-rods.

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71 Figure 3-11. Photoluminescence spectra from Gd2O3:Eu3+ micro-rods (black) and nano-rods (red). (a) PLE spectra for emission at 612 nm. (b) PL spectra excited at 280 nm. Decay time for (c) nano-rods and (d) micro-rods. Figure 3-12 shows the radioluminescence (RL) spectrum of calcined micro-rods irradiated by 40 keV X-ray. The spectrum is dominated by a peak at 612 nm originated in the 5D0 7F2 transition similar to the photoluminescence spectrum shown in Figure 3-11(b). In the particular case of Eu3+ doped Gd2O3, the forbidden 4f transitions become allowed by the Stark effect due to coupling of the Eu3+ energy levels with the crystal field of the host, leading to strong emission.

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72 Further, we analyzed the intensity ra tio of the emission lines at 589 nm (5D0 7F1) and 612 nm (5D0 7F2) that correspond to th e local symmetry insensitive ma gnetic dipole allowed transition and hypersensitive electric dipole transition, respectivel y. Differences in the intensity ratio of this particular pair of lines can be used to extrac t structural information in terms of the local symmetry of the Eu3+ ions. The results show the 5D0 7F2/5D0 7F1 intensity ratio to be ~9.6 and this indicates that the majority of Eu3+ ions are located in a site without inversion symmetry. Figure 3-12. Radioluminescen ce spectrum of calcined Gd2O3:Eu3+ micro-rods. 3.4 Conclusions Rod-like luminescent Gd2O3:Eu3+ phosphors were synthesized via a facile template-free sol-gel precipitation method at a low reaction temperature (~85 ) followed by calcining in air at 800 for 2 h. The sol-gel synthesis method produced hexagonal crystalline rods of Gd(OH)3:Eu3+ with diameters of 70 nm and lengths of either 200 nm or 1000 nm, designated nano-rods and micro-rods, respec tively. Nano-rods were produ ced when polyethylene glycol

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73 (PEG) was added as a capping agent. The rod geom etry and dimensions were maintained after calcining and the resultant c onversion of hexagonal Gd(OH)3:Eu3+ to cubic Gd2O3:Eu3+. HRTEM images show that the rod axis was paralle l to the hexagonal [001] crystal direction. The rod geometry was speculated to result from fast er growth along the caxis direction of the hexagonal Gd(OH)3 with slower growth rates in the [100] or [010] directions. The addition of PEG apparently led to slower growth along the [001] direction, and c onsequently led to the shorter nano-rods with a smaller aspect ratio (2). The photoluminescence (PL) and radioluminescence (RL) spectra from Gd2O3:Eu3+ were dominated by the 5D0-7F2 transitions from Eu3+ at 612 and 627nm. The decrease of PL intensity and lifetime from nano-rods compared to micro-rods was attributed to a larger concentration of non-radiative sites surrounding Eu3+ ions in nano-rods.

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74 CHAPTER 4 SYNTHESIS AND LUMINESCENT CHARACTER ISTIC OF EUROPIUM DOPANTS IN SILICA/GADOLINIUM OXIDE CORE/SHELL SCINTILLATION NANOPARTICLES 4.1 Introduction Luminescent rare-earth nanocry stals and colloidal quantum dot s doped with rare-earth ions are of interest due to their applications in va rious fields, e.g. displa ys, solid-state lighting, scintillator detectors, and biological diagnosis [104, 125-127] Gadolinium oxide exhibits good luminescent properties when doped with rare-earth ions. The characteristics of the luminescent and paramagnetic multifunctional properties lead to a number of possible applications such as neutron capture therapy targets and bimolecular detection [113, 114, 128]. Moreover, this oxide phosphor is a potential scintillati on detector material due to its short luminescent decay lifetime, high atomic number and density, low afterglow, large photon yield and lo w hygroscopicity [92, 96]. Traditionally, scintillation materials have been single crystal grown by demanding methods, i.e. Czochralski or Bridgman techniques. In th is work, a solution precipitation method was used to prepare SiO2:Eu3+/Gd2O3, SiO2/Gd2O3:Eu3+, SiO2/Gd2O3:Eu3+/SiO2 and SiO2/Gd2O3:Eu3+/Gd2O3 core/multi-shell scintillation nanopa rticles. With an ~13 nm thick Gd2O3:Eu3+ layer, the core/shell scint illating nanoparticles exhibited good luminescence. This aqueous solution processing offers a potentially cost-effective synthesis method for rare-earth doped scintillating material. In addition, SiO2 is a good host matrix for luminescent rare-earth ions due to its transparency, light weight and ease of producti on. Enhanced photoluminescence (PL) from Eu3+, Er3+ or Ce3+ ions in either SiO2 gel or thin film has been reported from ex-situ incorporation of discrete nanoparticles, such as CdS, Si, SnO2 and ZnO [68-70, 129, 130]. Nanocrystals may behave like phosphor sensitizers, harvesting the excitation energy from absorption of photons and subsequently transferring it to activators re sulting in increased PL. However, nanocrystals

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75 embedded by the ex-situ method must be within the required dipole interaction distance from the dopant ions in the host matrix in order to e ffectively transfer energy. In this work, monodispersed SiO2:Eu3+ nanoparticles were in-situ capped with a Gd2O3 layer resulting in a core/shell nanostructure with a four times larger quantum yield (QY). 4.2 Experimental 4.2.1 Synthesis of Core/Shell and Core/Multi-Shell Nanoparticles Silica nanoparticles were fabricated via the Stober method [131], in which 10 mL of concentrated ammonium hydroxide (NH4OH) was mixed with 200 mL of ethanol and 40 mL of deionized (DI) water and stirred thoroughly. Then, 20 mL of tetr aethyl orthosilicate (TEOS) solution was added and stirred vigorously for 40 mi n. No visible change occurred at first but after a few minutes, the solution became hazy and then opaque white as SiO2 particles grew large enough to scatter light. Mono-dispersed 220~230 nm diameter spherical SiO2 nanoparticles were obtained and used as cores for further coating. To produce Eu3+ doped SiO2 cores, 0.008 mol/L of europium nitrate was added before the add ition of TEOS. Samples were collected by washing and centrifuging with DI water and ethanol several times before further procedures. A gadolinium oxide shell was grown on the SiO2 cores by a solution precipitation method. A solution of 0.4 g/100 mL of SiO2 nanoparticles in DI water was prepared and sonicated to achieve a homogeneous suspension, followed by a dditions of 0.04 mol/L of gadolinium nitrate and 0.008 mol/L of europium nitrate mixed with 1 mol/L urea at 82~85 to yield a transparent stirred solution. After vigorous stirring for 2~4 h, the white preci pitate was collected by washing and centrifuging with DI water severa l times, followed by oven-drying at 80 For multi-shell nanoparticles SiO2 coating, previous synthesized 1.8 g SiO2/Gd2O3:Eu3+ core/shell nanoparticles were mixed with 60 mL of DI water and 250 mL of ethanol, followed by addition of 5 mL of NH4OH and 1 mL of TEOS added dr op-wise into the solution a nd reacted for 1~3 h. For

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76 SiO2/Gd2O3:Eu3+/Gd2O3 core/multi-shell nanoparticles, similar procedures were followed as for previous precipitation steps except that no Eu3+ dopant ions were added, and the reaction time was 90 min. The products were oven-dried at 80 for 2 h, followed by calcination at 800 in air for 2 h. 4.2.2 Characterization As-prepared and calcined samples were characterized by X-ray diffraction (XRD) (Philips APD 3720) with Cu K radiation source (=1.5418 ). The XRD pattern was collected from dried powder samples in the step scan (0.02) mode over a 2 range of 20-70. The morphology and size of core/shell nanoparticle s were determined with high resolution transmission electron microscopy (HR-TEM) and with field emission scanning electron microscopy (FE-SEM). The TEM samples were prepared by drop-casting samp les dispersed in ethanol onto a carbon-coated holey copper grid followed by drying at room temperature. The transmission spectra of the samples and standards solution were measured using a UV/Vis Perkin-Elmer Lambda 800. Photoluminescence (PL) and photoluminescence ex citation (PLE) spectra were measured at room temperature using a JASCO FP-6500/6600 resear ch-grade fluorescence spectrometer with a 150 W xenon lamp. Quantum yields were measur ed using an integrating sphere and the samples were prepared by drop-casting (30 mg of calcined nanoparticles in 500 L of 4 % polymethyl methacrylate (PMMA) dispersed in a chloroform/benzene mixture) onto quartz substrates. 4.3 Results and Discussion 4.3.1 Structural Analysis The morphology and size of as-prepared spherical SiO2 nanoparticles are shown in Figure 4-1(a). Based on SE M micrographs, SiO2 nanoparticles are mono-disper sed with a diameter of ~220 nm. A thin ~13 nm Gd2O3 crust was added onto the SiO2 cores by the urea precipitation

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77 process, as shown by the TEM micrographs in Figures 4-1(b) and (c). In this method, urea is the precipitation agent and Gd and Eu metal ions are precipitated on the co re surface with more growth at longer reaction time. The shell thickn ess is a function of the urea concentration, reaction time and temperature. It is clear that a uniform SiO2/Gd2O3:Eu3+ core/shell nanostructure was obtained and the outer dark and inner light regi ons correspond to the Gd2O3 shell and SiO2 core, respectively, due to the difference in atomic number. Figure 4-1(d) shows the core/shell nanoparticles af ter calcining in air at 800 for 2 h. Note that well-resolved crystalline lattice fringes were observed in the outer dark region indicating that the Gd2O3 shell is crystalline. The inset of Figure 4-1(d) s hows the bright red PL from calcined SiO2/Gd2O3:Eu3+ core/shell nanoparticles under room light and ultra-violet (UV) excitation. Figure 4-1. (a) SEM microgr aph of mono-dispersed SiO2 nanoparticles synthe sized by the Stober method. TEM micrographs of as-prepared SiO2/Gd2O3:Eu3+ core/shell nanoparticles at (b) low magnification and (c) high magnification. (d) TEM micrograph of SiO2/Gd2O3:Eu3+ core/shell nanoparticles af ter being calcined at 800 for 2 h in air. The inset shows the luminescence under r oom light or UV irradiation from a Hg discharge lamp.

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78 XRD data in Figure 4-2(a) indicate that the SiO2 nanoparticles are amorphous while asprepared SiO2/Gd2O3:Eu3+ core/shell nanoparticles show extr emely broad peaks, as shown in Figure 4-2(b), cons istent with nano-grain size cubic Gd2O3. After the samples were calcined at 800 for 2 h, a well-crystallized Gd2O3 phase was obtained which matched cubic Gd2O3 (JCPDS: #43-1014; Figure 4-2(d)). This core/she ll nanostructure has the advantages of easy synthesis, monodispersity and good control of core size. Further, uniform spherical shells can be coated on SiO2 core resulting in a mono-disper sed core/shell nanostructure. 1020304050607080 Cubic Gd2O3 (#43-1014) 2 theta SiO2/Gd2O3:Eu3+ Calcined@8000CIntensity (Counts) SiO2/Gd2O3:Eu3+ (As-prepared) Pure SiO2(a) (b) (c) (d)1020304050607080 Cubic Gd2O3 (#43-1014) 2 theta SiO2/Gd2O3:Eu3+ Calcined@8000CIntensity (Counts) SiO2/Gd2O3:Eu3+ (As-prepared) Pure SiO2(a) (b) (c) (d) Figure 4-2. XRD spectra obtained from (a) pure silica nanoparticles, (b) as-prepared SiO2/Gd2O3:Eu3+ core/shell nanoparticles, and (c) calcined SiO2/Gd2O3:Eu3+ core/shell nanoparticles (800 for 2 h in air). (d) Standard diffraction peaks from JCPDS (#43-1014) for cubic Gd2O3. In order to study the effects of surface pa ssivation on PL, core/she ll/shell scintillation nanoparticles were synthesized, e.g. SiO2/Gd2O3:Eu3+/SiO2 and SiO2/Gd2O3:Eu3+/Gd2O3. Figure

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79 4-3(a) shows a TEM microgr aph of as-prepared SiO2/Gd2O3:Eu3+/SiO2 nanoparticles. Three different regions can be observe d showing the core/shell/shell na nostructures as indicated by the arrows. A uniform ~10 nm thick SiO2 shell is observed on the doped Gd2O3:Eu3+ shell on the SiO2 core. As shown in Figure 4-3(b), good coverage by the Gd2O3:Eu3+ and SiO2 shells was observed even after the sample was calcined at 800 for 2 h. 20nm(b) 10nm SiO2 Gd2O3:Eu3+SiO2 (a) 20nm(b) 20nm 20nm(b) 10nm SiO2 Gd2O3:Eu3+SiO2 10nm SiO2 Gd2O3:Eu3+SiO2 (a) Figure 4-3. TEM micrographs of core/shell /shell nanostructure of (a) as-prepared SiO2/Gd2O3:Eu3+/SiO2 nanoparticles and (b ) calcined SiO2/Gd2O3:Eu3+/SiO2 nanoparticles. The optical and structural propertie s of as-prepared and calcined SiO2 nanoparticles were studied. Figure 4-4(a) shows the transmi ssion spectra in which the as-prepared SiO2 absorbed strongly below ~330 nm, wh ile after calcining at 800 for 2 h they show good transparency in the UV region for >250 nm. These data suggest that SiO2 shell absorbs ~15 % of the UV necessary to excite PL from the Eu3+. The XRD data in Figure 4-4(b) shows that both asprepared and calcined SiO2 samples are amorphous.

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80 300400500600700 0 20 40 60 80 100 Transmission (%)Wavelength (nm) As-prepared SiO2 Calcined SiO210203040506070 As-prepared SiO2 Calcined SiO2 Intensity (a.u.)2 theta (Degree)(a) (b)300400500600700 0 20 40 60 80 100 Transmission (%)Wavelength (nm) As-prepared SiO2 Calcined SiO210203040506070 As-prepared SiO2 Calcined SiO2 Intensity (a.u.)2 theta (Degree)(a) (b) Figure 4-4. (a) Optical transmission spectra obtained from as-prepared and calcined SiO2 nanoparticle layer. (b) XRD spectra of as-prepared and calcined SiO2 nanoparticle layer. 4.3.2 Thin Film Quantum Yield Measurement Room-temperature thin-film photoluminescen ce was measured with a spectrophotometer (FP-6500, Jasco, Inc.) equipped with a DC-powered 150W Xenon lamp source and a photomultiplier tube (PMT) detector. A quantum yi eld measurement system for solid thin film samples acquires the light with an integrating sphere (60 mm diameter; BaSO4 coating; Spectralon reflectance standards). The quantum yi eld (QY) of core/shell and multi-shell scintillation nanoparticles is defi ned to be the fractional or per centage ratio of the number of emitted photons to the number of absorbed photons, i.e. S2 divided by (S0-S1) as illustrated in Figure 4-5. Here S0 is the number of incident photons measured using a reflective standard, S1 and S2 are the number of incident photons not absorbed and luminous photons emitted by thin film samples, respectively [132, 133].

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81 Sample S2: Number of luminous photons S1: Number of photons not absorbed S0: Number of irradiated photons Blank Quantum Yield (%) # of Emitted Photons # of Absorbed Photons = S0-S1S2= Sample S2: Number of luminous photons S1: Number of photons not absorbed S0: Number of irradiated photons Blank Quantum Yield (%) # of Emitted Photons # of Absorbed Photons = S0-S1S2= Figure 4-5. Schematic diagram of thin film quantum yield (QY) measurement. 4.3.3 Luminescent Properties of Core/Multi-Shell Nanoparticles Figure 4-6(a) shows PLE spectra for SiO2/Gd2O3:Eu3+ core/shell and SiO2/Gd2O3:Eu3+/SiO2 core/shell/shell nanopartic les. The broad PLE peak at <300 nm is from the oxygen to europium charge transfer band (CTB). The gadolinium 8S6I7/2 transition around 280 nm is observed and overlaps the wavelength range of the charge -transfer band of Eu3+ [124, 134, 135]. The narrow peaks from 350 to 550 nm ar e the europium self-excitation peaks, and they do not change in waveleng th with the addition of a SiO2 shell. However, the oxygen to europium CTB peak is much less intense for SiO2/Gd2O3:Eu3+/SiO2 as compared with SiO2/Gd2O3:Eu3+ samples. As shown in Figure 4-6(b), due to the weak CTB, SiO2/Gd2O3:Eu3+/SiO2 samples show lower PL intensity as compared to SiO2/Gd2O3:Eu3+ samples, even though both PL spectra are dominated by the 5D0-7F2 transitions with peaks at 609 and 622nm. Decreased PL emission from SiO2/Gd2O3:Eu3+/SiO2 core/shell/shell nanoparticles is attributed to the ~15 % UV absorption at 280 nm along with a higher density of interface quenching sites introduced by the amorphous SiO2 shell.

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82 (a)575600625650675700725 0 50 100 150 SiO2/Gd2O3:Eu3+ SiO2/Gd2O3:Eu3+/ SiO2PL Intensity (a.u.)Wavelength (nm)Ex: 280 nm 5D0 -7F1 5D0 -7F2 5D0 -7F3 5D0 -7F4(b)250300350400450500550 PLE Intensity (a.u.)Wavelength (nm) SiO2/Gd2O3:Eu3+ SiO2/Gd2O3:Eu3+/ SiO2Em: 609 nm CTB(O2-Eu3+) Eu3+Gd(8S-6I7/2)(a)575600625650675700725 0 50 100 150 SiO2/Gd2O3:Eu3+ SiO2/Gd2O3:Eu3+/ SiO2PL Intensity (a.u.)Wavelength (nm)Ex: 280 nm 5D0 -7F1 5D0 -7F2 5D0 -7F3 5D0 -7F4(b)575600625650675700725 0 50 100 150 SiO2/Gd2O3:Eu3+ SiO2/Gd2O3:Eu3+/ SiO2PL Intensity (a.u.)Wavelength (nm)Ex: 280 nm 5D0 -7F1 5D0 -7F2 5D0 -7F3 5D0 -7F4575600625650675700725 0 50 100 150 SiO2/Gd2O3:Eu3+ SiO2/Gd2O3:Eu3+/ SiO2PL Intensity (a.u.)Wavelength (nm)Ex: 280 nm 5D0 -7F1 5D0 -7F2 5D0 -7F3 5D0 -7F4(b)250300350400450500550 PLE Intensity (a.u.)Wavelength (nm) SiO2/Gd2O3:Eu3+ SiO2/Gd2O3:Eu3+/ SiO2Em: 609 nm CTB(O2-Eu3+) Eu3+Gd(8S-6I7/2) Figure 4-6. Photoluminescence spectra from SiO2/Gd2O3:Eu3+ core/shell and SiO2/Gd2O3:Eu3+/SiO2 core/shell/shell nanopa rticles. (a) PLE spectra for emission at 609 nm. (b) PL spectra excited at 280 nm. For excitation at 280 nm, the QY was ~25 %, 17 % and 32 % from SiO2/Gd2O3:Eu3+, SiO2/Gd2O3:Eu3+/SiO2 and SiO2/Gd2O3:Eu3+/Gd2O3, respectively, as shown in Table 4-1. The increased QY from adding an undoped Gd2O3 second shell may result from reduced surface inhomogenities and passivation of surface que nching sites or defects along with energy harvesting and subsequent transfer from the undoped second shell to Eu3+ in the doped Gd2O3 first shell. Brighter PL from core/undoped homo-shell nanostructures has been reported [62, 136]. The reduced PL intensity from SiO2/Gd2O3:Eu3+/SiO2 structures may result from increased interfacial quenching sites introduced by the amorphous SiO2 second shell. Table 4-1. Quantum yields of core/shell and co re/shell/shell nanoparticles excited at 280 nm. Excitation Wavelength SiO2/Gd2O3:Eu3+ SiO2/Gd2O3:Eu3+/SiO2 SiO2/Gd2O3:Eu3+/Gd2O3 280 nm 25.0 16.5 32.2 4.3.4 Luminescent Properties of SiO2:Eu3+ and SiO2:Eu3+/Gd2O3 Nanoparticles To evaluate the effects of surface modifications of SiO2 on photoluminescence, a SiO2:Eu3+ core without and with an undoped Gd2O3 shell were synthesized and calcined at 800

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83 for 2 h in air. Figure 4-7(a) show s the PLE and PL spectra of the SiO2:Eu3+ and SiO2:Eu3+/Gd2O3 nanostructures. The data show that na noparticles consisting of a core capped with a Gd2O3 shell have much brighter PL emission at 609 and 622 nm when excited by 278 nm UV, as compared to the bare core. The QYs of SiO2:Eu3+ cores without and with a Gd2O3 shell, excited at 278 nm, are shown in Figure 4-7(b) as a function of Gd2O3 coating time. The insets are photos of nanoparticle powders irradiated with UV from a mercury discharge lamp. The QY of bare SiO2:Eu3+ nanoparticles (0 h) is ~2.2 %, but is increased by growth of a Gd2O3 shell and reaches a maximum of ~8.5 % after a reaction time of 3 h. This enhancement is attributed to the undoped crystalline Gd2O3 layer acting as a sensitizer, harvesting energy at 278 nm and subsequently transferring it to the Eu3+ in the SiO2 core. As discussed above, the shell may also passivate non-radiative surface defect states on the SiO2 core. 01234 0 1 2 3 4 5 6 7 8 9 10 Quantum Yield (%)Coating Time (Hour)01234 SiO2:Eu3+/Gd2O3 SiO2:Eu3+ (b) (a) 300400500600700 PLE Intensity (a.u.)Wavelength (nm) PL Intensity (a.u.) SiO2:Eu3+ SiO2:Eu3+/ Gd2O3Gd(8S-6I7/2) 01234 0 1 2 3 4 5 6 7 8 9 10 Quantum Yield (%)Coating Time (Hour)01234 SiO2:Eu3+/Gd2O3 SiO2:Eu3+ 01234 0 1 2 3 4 5 6 7 8 9 10 Quantum Yield (%)Coating Time (Hour)01234 SiO2:Eu3+/Gd2O3 SiO2:Eu3+ (b) (a) 300400500600700 PLE Intensity (a.u.)Wavelength (nm) PL Intensity (a.u.) SiO2:Eu3+ SiO2:Eu3+/ Gd2O3Gd(8S-6I7/2) Figure 4-7. (a) Photolumin escence spectra from SiO2:Eu3+ and SiO2:Eu3+/Gd2O3 core/shell nanoparticles with PLE spectra for emission at 609 nm and PL spectra excited at 278 nm. (b) Quantum yield of SiO2:Eu3+ and SiO2:Eu3+/Gd2O3 as a function of Gd2O3 reaction time with the inse ts showing photos of PL excited by UV irradiation.

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84 4.3.5 Photoluminescent Enhancement Mechanism We speculate that energy transfer leads to the enhanced PL from SiO2/Gd2O3:Eu3+/Gd2O3 and SiO2:Eu3+/Gd2O3 nanoparticles, possibly by the process illustrated in Figure 4-8. Sensitizers and activators are denoted as S and A, respectively, where an asterisk indi cates an excited state. The sharp line around 280 nm in Fi gure 4-7(a) corresponds to the Gd3+ 8S6I transition which overlaps the wavelength range of the charge-transfer band of Eu3+. Under excitation around 280 nm, the Gd3+ ions are excited into the 6Ij states, followed by non-radiative energy transfer to a neighboring Eu3+ ion in the SiO2 core or Gd2O3 shell, resulting in enhanced PL. The use of Gd3+ as a sensitizer for en ergy transfer to Eu3+ ions has been reported [ 137, 138]. This energy transfer would be inhibited for Gd2O3 thicknesses larger than the cr itical distance for dipole-dipole interactions, which is reported to be a fe w nanometers in gadolinium compound [139-141]. Therefore, thick undoped Gd2O3 shells would absorb UV leading to lower PL intensities. Similar variations in PL intensity have been reporte d for doped nanorods or na noparticles capped with undoped shells [23, 62, 142, 143]. Gd3+ Energy Transfer Eu3+ in SiO2 or G2O3 7Fj 5D0 5D2 Emission A* S* S A 8S7/2 6Pj 6Ij 6Dj Energy (103cm-1)0 10 20 30 40 50 Gd3+ Energy Transfer Eu3+ in SiO2 or G2O3 7Fj 5D0 5D2 Emission A* S* S A 8S7/2 6Pj 6Ij 6Dj Energy (103cm-1)0 10 20 30 40 50 Energy (103cm-1)0 10 20 30 40 50 Figure 4-8. Schematic illustration of po ssible charge transfer paths in SiO2:Eu3+/Gd2O3 core/shell and SiO2/Gd2O3:Eu3+/Gd2O3 nanoparticles.

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85 4.4 Conclusions Core/shell nanostructures with mono-dispersed 220 nm SiO2 cores and ~13 nm Gd2O3 were prepared by a solution precipitation me thod. Photoluminescence (PL) was achieved by doping either the SiO2 core or Gd2O3 shell with Eu3+ ions. A thin Gd2O3:Eu3+ shell exhibited good PL with a quantum yield (QY) up to 25 %. In addition, SiO2/Gd2O3:Eu3+/X core/shell/shell nanostructures were synthesized with X being undoped SiO2 or Gd2O3. The amorphous SiO2 second shell decreased the PL in tensity, while the crystalline Gd2O3 second shell increased the intensity from the inner Gd2O3:Eu3+ shell. An undoped Gd2O3 shell on a SiO2:Eu3+ core increased the QY from ~2.2 % for the bare core to a maximum of ~8.5 %. The increased PL intensity was attributed to the crystalline Gd2O3 layer serving as a sensit izer with energy transfer to the Eu3+ either in the Gd2O3:Eu3+ shell or SiO2:Eu3+ core along with the passivation of nonradiative surface defect states.

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86 CHAPTER 5 CORE/SHELL COMPOSITE OF SELF-A SSEMBLED HIERARCHICAL BISMUTH OXIDE/EUROPIUM-DOPED GADOLINIUM OXIDE 5.1 Introduction A number of applications of bismuth oxide have been widely studi ed, including optical coating, gas sensor, solid-state electrolytes, micr oelectronics and electroceramic, due to its optoelectronic, catalyst, photoluminescent and ion-co nducting properties [144147]. Bismuth oxide is a complex material with four ma in polymorphs which are denoted by (monoclinic), (tetragonal), (body centered cubic) and -Bi2O3 (face centered cubic) along with a new modification, bismuth oxide [148]. Each polymorph is ch aracterized with distinct structural, optical and electrical properties. The and phases are the stable phases at room temperature and high temperature (between 729 and the melting point at 824 ), respectively. The others are high temperature metastable phases and usua lly transforms into the stable monoclinic Bi2O3 [149, 150]. Among them, -Bi2O3 has the best known oxide ion conductivity which makes it appealing for fuel cell and sens or applications, but its narrow stable temperature range of 729~824 along with cracking and dete rioration during phase transi tions limits its use. The addition of rare-earth dopants can stabilize th e cubic phase, but the c onductivity drops by two orders of magnitude and stability remains an i ssue. Several study have been reported on the synthesis of -Bi2O3 in the form of spheres or one-dime nsional structure using techniques such as chemical bath deposition (CBD), electros pinning, hydrothermal synthesis, flame spray pyrolysis (FSP), and microemulsion methods [151-154]. In addition, -Bi2O3 has been synthesized by epitaxial electr odeposition onto cubic Au single crystal substrates or by introducing VO3 into the reaction system [155]. However, there are no reports of -phase Bi2O3 having been synthesized at low temperatures in short reaction times without subsequent heat treatment. Several methods have be en reported for the synthesis of Bi2O3 nanoparticles,

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87 nanowires, nanotube, nanofibers and nanohooks [153, 156, 157], but there are only a few reports on the synthesis of three-dimensional (3D) hier archical structures. To date, strategies for synthesis of these materials with higher dime nsion structures under template-free and costeffective conditions remain a challenge. In this work, 3D flower-like -Bi2O3, composed of 2D building blocks, was synthesized via a facile solution precipitation method using low reaction temperatures and short times with the aid of polyethylene glycol-8000 (PEG-8000). Structural, photoluminescent and conductivity properties of hierarchical -Bi2O3 are reported. 5.2 Experimental 5.2.1 Synthesis of Rod-Like Bi2O3 and Hierarchical Flower-Like Bi2O3 All chemicals were analytical grade and used as received wi thout further treatment. The typical solution precipita tion procedure was begun by dissolving 1.94 g of Bi(NO3)35H2O in 20 mL of HNO3 (1M) to form a clear solution, followe d by adding 200 mL of deionized water. The precursor mixture was heated to 85 and different concentrations of PEG-8000 (designated as C and defined as the volume ratio of PEG-8000 to the total precursor solution) was added to the solution with vigorously stirring until a hom ogeneous clear solution was again obtained. Then 60 mL of NaOH (4M) solution was added to the precursor mixture w ith continuous stirring to generate a colloidal solution. After various reaction times, the prec ipitate was collected by washing with DI water and centrifuging several times, followed by oven-drying at 80 for 2 h. 5.2.2 Synthesis of Bi2O3/Gd2O3:Eu3+ Core/Shell Composite Luminescent gadolinium oxide shell was coated onto the Bi2O3 cores using a sol-gel method. A concentration of 0.4g/100mL of Bi2O3 was prepared to have a homogeneous suspension, followed by additions of 0.05 mol/L gadolinium nitrates and 0.008 mol/L europium nitrates mixed with 1 mol/L urea to yield a precursor solution that was stirred at 82~85 After being vigorously stirring for 2 h, precipitates were collected by washing and centrifuging with DI

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88 water several times. The products were oven-dried at 80 for 2 h, followed by the calcinations at 500 in air for 2 h. 5.2.3 Characterization As-prepared and calcined (600 for 2 h in air) samples were characterized by X-ray diffraction (XRD) (Philips APD 3720) with Cu K radiation (=1.5418 ). The XRD pattern was collected from dried powder samples in a step scan (0.02 per step) mode over a 2 scan range of 20-70. The morphology and size of Bi2O3 were determined with high resolution transmission electron microscopy (HR-TEM; JE OL 2010F) and with field emission scanning electron microscopy (FE-SEM; JEOL 6335F). Photoluminescence (PL) and photoluminescence excitation (PLE) spectra were measured at room temperature us ing a JASCO FP-6500/6600 spectrophotometer with a 150 W xenon lamp. Conductivity was determined by two-point probe electrochemical impedance spec troscopy (EIS) using a Solartron 1260 over the frequency range of 1 Hz to 1 MHz and between 300~600 in air. 5.3 Results and Discussion 5.3.1 Rod-Like and Hierarchical Flower-Like Bi2O3 5.3.1.1 Capping agent effect-crystal and morphology study To determine the effects of a surface cappi ng agent on crystal growth and morphology, different volume concentrations of PEG-8000 were added to the precursor solution with other experimental parameters held constant. SE M micrographs show that the morphology of Bi2O3 was changed significantly by increased PEG con centration (Figure 5-1) The morphology was micro-rods when either zero or lo w concentrations of PEG-8000 (C < 0.05) were used, as shown in Figure 5-1(a). With an increased volume co ncentration of PEG-8000 (C = 0.1 or 0.2), the morphology evolved from micro-rods into flow er-like crystals (Fi gure 5-1(b) and (c), respectively). Based on the Gibbs-Curie-Wulff theorem, the morphology can be explained by the

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89 relative growth kinetics of different facets of the crystal [3, 48]. Introduction of a capping agent into the reaction can tailor the growth kinetics by modulating the relative free energies of various crystallographic facets. A similar growth mechanism was proposed to explain the growth of ZnO nanocones and FeCo nanocubes in the presence of PEG [120, 158]. Figure 5-1. SEM images of morphology evolution as a function of PEG-8000 volume fraction, C = (a) 0.05, (b) 0.1, and (c) 0.2 (r eaction temperature = 85 time = 45 min). Figure 5-2 shows the XRD spectra as a f unction of PEG-8000 volume concentration, which indicates that the two distin ct morphologies are associated w ith different crys tal structures. The spectrum in Figure 5-2(a) indicates that the micro-rods formed at C < 0.05 are monoclinic phase (JCPDS card #71-2274; Figure 5-2(d)). Howe ver, when the volume concentration of PEG8000 was increased to C = 0.1 and then to 0.2, the diffraction patterns in Figure 5-2(b) and (c) were from both the monoclinic and body centered cubic phases, but was dominant for the flower-like morphology from C = 0.2 (JCPDS card #81-0563; Figure 5-2(e)). Previously -Bi2O3 was reported to form at temperatures of 640 or higher [149].

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90 203040506070 203040506070 203040506070 203040506070 203040506070 Intensity (a.u.)2 theta (Degree) Cubic (#81-0563) Monoclinic (#71-2274) PEG: C=0.2 PEG: C=0.1 PEG: C=0.05 (a) (b) (c) (d) (e)203040506070 203040506070 203040506070 203040506070 203040506070 Intensity (a.u.)2 theta (Degree) Cubic (#81-0563) Monoclinic (#71-2274) PEG: C=0.2 PEG: C=0.1 PEG: C=0.05 (a) (b) (c) (d) (e) Figure 5-2. X-ray diffraction (XRD) pattern from bismuth oxide as a function of PEG-8000 volume fraction, C = (a) 0.05, (b) 0.1, a nd (c) 0.2 (reaction temperature = 85 time = 45 min). (d) Diffraction peaks from JCPDS for monoclinic -bismuth oxide. (e) Diffraction peaks from JCPDS for cubic -bismuth oxide. 5.3.1.2 Time-dependent growth analysis of hierarchical flower-like Bi2O3 To investigate the morphology evolution and formation mechanisms of the 3D selfassembled hierarchical architecture, a growthtime analysis was performed with XRD and SEM. The morphology and size of the 3D flower-like bi smuth oxide structure as a function of reaction time from 1 min to 45 min (reaction temperature = 85 C = 0.2) are shown by the SEM micrographs in Figure 5-3. The morphology evol ved from ~60 nm nano-spheres after 1 min reaction time (Figure 5-3(a)), to agglomerated sub-micron clusters after 10 min (Figure 5-3(b)), to predominantly 3D flowers af ter 45 min (Figure 5-3(c)).

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91 Figure 5-3. Evolution of the mor phology of bismuth oxide as a func tion of reaction times of (a) 1 min, (b) 10 min and (c) 45 min (C = 0.2, reaction temperature = 85 ). Figure 5-4 shows the XRD spectra from as-pre pared samples for reaction times from 1 min to 45 min. Figure 5-4(a) shows a broad peak with a large full width at half maximum (FWHM) that is attributed to amorphous nano-size grains. The XRD pattern s for reaction times of 10 min and 45 min are shown in Figure 5-4(b) and (c), respectively. Two XRD spectra are shown in Figure 5-4(c), with the black spectrum being from a sample grown for 45 min, and the red spectrum from a sample grown for 45 min th en calcined at 600 in air for 2 h. Both the asprepared (black) and calcined samples (red) sh ow a strong crystalline pattern matched by the Bi12.8O19.2 phase (i.e. -Bi2O3) which is composed of umbrella-like [BiO3] groups and void tetrahedral (JCPDS: 81-0563; Figure 5-4(d)). In addition, the morphol ogy was not changed by calcining as shown in Figure 5-5 (a)-(c).

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92 203040506070 203040506070 203040506070 203040506070 Intensity (a.u.)2 theta (Degree)Cubic (#81-0563) 45 min 10 min 1 min(a) (c) (d) (b)203040506070 203040506070 203040506070 203040506070 Intensity (a.u.)2 theta (Degree)Cubic (#81-0563) 45 min 10 min 1 min(a) (c) (d) (b) Figure 5-4. XRD pattern from bismuth oxide as a function of reac tion times of (a) 1 min, (b) 10 min and (c) 45 min (black: as-grown; red: calcined at 600 for 2 h in air), (C = 0.2, reaction temperature = 85 ). (d) Diffraction peaks from JCPDS for cubic -bismuth oxide. Figure 5-5. SEM micrographs at different magnifications of flower-like Bi2O3 calcined at 600 for 2 h in air showing that calcining does not change the morphology.

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93 5.3.1.3 Structural and photoluminescent prop erties of hierarchical flower-like Bi2O3 These 3D bismuth flowers have high hierarchy with petals which are composed of selfassembled nano-triangular and pyramid structures as shown in Figure 5-6(a). The typical size of these individual nano-triangles are 600~800 nm. The nano-tria ngles form a petal which originates from a single center to exhibit the fl ower-like geometry. The size of an assembled flower structure is ~10 m. Figure 5-6(b) shows a high resolution TEM image of a bismuth oxide nano-triangle revealing well-resolv ed lattice fringes with an in terplanar spacing of 0.256 nm. In addition, the selected area elec tron diffraction (SAED) pattern in the inset of Figure 5-6(b) shows single crystal spots from the cubic -phase, consistent with the XRD patterns. Figure 5-6. (a) SEM micrograph of an individual flower of bi smuth oxide. (b) TEM micrograph of the lattice fringes from an individual nano-triangle with its SAED pattern shown in the inset. Figure 5-7 shows a room-temperature photoluminescence (PL) spectrum from -phase asprepared flower-like Bi2O3 (excited at 230 nm). The PL spectrum shows broad emission (~400650 nm) with a peak at 465 nm due to Bi3+ luminescence from the 3P0 and 3P1 excited states to the 1S0 ground state. The green peak at ~562 nm is from an impurity trap associated with oxygen vacancies interacting with inte rfacial bismuth vacancies [157] The PL transition observed

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94 depends on the charge on the bismuth ion, and th e blue and green emissions are consistent with trivalent bismuth ions. 350400450500550600650700 PL Intensity (a.u.)Wavelength (nm) PL spectrum of flower-like Bi2O3 Figure 5-7. Photoluminescence spectru m of as-synthesized flower-like -bismuth oxide for excitation at 230 nm. 5.3.1.4 Possible formation mechanism A schematic illustration of a possible formati on process leading to growth of 3D selfassembled flower-like and rod-like bismuth oxide cr ystals is shown in Figure 5-8. In general, the crystalline phase, particle size and morphology of a solid depend strongly on the competition between nucleation and growth, th e rates of which are determined by the chemical potentials in the precursor solution. During the initial reac tion, bismuth oxide proba bly nucleated as the reagent acid was poured into the pr ecursor solution, and grew to nano-size particles in < 1 min. With continued reaction, nano-size particles were oriented and attached to other nano-particles, leading to agglomerated nano-particles which se rve as crystal seeds to grow the flower-like structure. PEG-8000 is an amphipathic non-i onic surface active agent from oxyethylene polymerization, and therefore a PEG-OH bond should form in aqueous solution and chelate

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95 bismuth ions, which would lead to highly mobile molecules with a larg e exclusion volume. The PEG chains result in the morphology of the initially formed bismuth oxide clusters. Moreover, due to the large amount of high molecular wei ght PEG-8000, the precursor solution is viscous leading to the agglomeration of nano-size bismut h oxide clusters to form submicron-triangular platelets. Finally, PEG-8000 leads to bridging flocculation which bi nds triangular platelets into individual petal-like structur es, resulting in flower-like -Bi2O3. Figure 5-8. Schematic illustration of a possible formation process for 3D flower-like bismuth oxide. 5.3.2 Bi2O3/Gd2O3:Eu3+ Core/Shell Composite 5.3.2.1 Structural analysis Figure 5-9(a) and (b) show the SEM micrographs of flower-like bismuth oxide before and after capping of europium doped gadolinium oxide, respectively. Clearly, the individual composite/cluster in Figure 5-9( b) is larger than bismuth oxide flowers in Figure 5-9(a). The flower-like feature of bismut h oxide disappeared and was en compassed by gadolinium oxide materials. As shown in Figure 5-9(c), th e higher magnification SEM micrograph shows

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96 individual core/shell scintillati on composite. It can be seen th at the bismuth oxide petal was covered with gadolinium oxide nanocrystals precipit ated on the petals duri ng the sol-gel process. This scintillation core/shell composite synthesized by th e facile sol-gel method has the advantages such as high stopping power, ease of synthesis, stable and non-hydroscopic properties. Figure 5-9. SEM micrographs of (a) self-asse mbled flower-like bismuth oxide, (b) bismuth oxide/europium doped gadolinium oxide core/shell scintillation composite and (c) high magnification from the indi vidual core/shell composite. 5.3.2.2 Photoluminescence analysis Photoluminescent excitation (PLE) and photoluminescent emission (PL) spectra for asprepared and calcined samples are shown in Figur e 5-10(a) and 5-10(b), re spectively. In Figure 5-10(a), the broad peak at <300 nm is from the oxygen to europium charge transfer band (CTB). The sharp peaks between 350 to 550 nm are the europium self-excitation peaks, which are generally less dependent of the lattice structure or crystal size of the matrix. The oxygen to europium CTB is much le ss intense for as-prepared samples as compared with calcined samples. Due to the strong CTB, calcined samples show higher PL emission from 5D0 to 7Fj europium transitions as shown in Figure 5-10(b). Shown in the inset of Figure 5-10(b ) are the photos of asprepared and calcined samples under room lighting and ultra-violet (UV) irradiation excited by a mercury-discharge lamp. This demonstrates th at the scintillation co mposite exhibits good luminescence upon UV irradiation.

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97 Figure 5-10. Photoluminescent analysis from bismuth oxide/gadolinium oxide core/shell scintillation composite. (a) PLE spectra observed at 612 nm. (b) PL spectra excited at 280 nm. The inset shows the photos under room lighting and UV irradiation. 5.4 Conclusions Unique hierarchical 3D self-assembled flower-like -Bi2O3 was produced in the presence of PEG-8000 via a facile solution precipitation process without any templates. This is the first report of the growth of body centered cubic -phase Bi2O3 at such a low temperature (85 ) and a short time (45 min). By varying the reaction time and PEG-8000 capping agent concentration, the evolution of morphology and crystal phase of bismuth oxide was investigated. The present data show that the -phase Bi2O3 formed at 85 is stable up to at least 600 The photoluminescent properties and possible growth mechanisms were discussed. Further deposition of europium doped gadolinium oxide on this hier archical bismuth oxide leads to core/shell scintillation composite. The composite exhi bits good luminescent pr operties upon calcination indicating a good coverage of gadolinium oxide laye r and the potential use in the scintillation detector.

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98 CHAPTER 6 SYNTHESIS AND CHARACTERIZATION OF SELF-ACTIVATED BISMUTH GERMANIUM OXIDE SCINTILLATION MATERIALS 6.1 Introduction Detection of X-rays, -rays, and charged particles i nvolves the energy deposited by electrons or charged particles gene rated by the incident radiation. A scintillator is a material with the ability to absorb the ionizi ng radiation and convert a fraction of the energy into visible or ultraviolet photons that are detected by a photodiode or photomultiplier (PMT). A good scintillation material should be luminescent wi th a short decay time, high atomic number and density, low afterglow, large light output a nd low hygroscopicity [96, 98, 159]. Scintillation materials are important to a number of applicatio ns, e.g. high energy physic s, security inspection stations, nuclear detection, i ndustrial control, medical imagi ng and petroleum exploration. Bismuth germanate (Bi4Ge3O12-BGO) is one of the most popul ar single crystal scintillation material for positron emission tomography (PET) medical imaging systems [94, 160]. BGO has good scintillator properties such as high density (7.13 g/cm3), high effective atomic number (Z=524), relatively short decay time (300 ns ), good energy resolution (10-15% FWHM), radiation hardness, high light output (93 photons/MeV), low afterglow (0.005% after 3 ms), rugged (no cleavage plane) and good sensitiv ity to a variety of particles (UV, and ) [161, 162]. Generally, the scintilla tors are single crystals with lim ited dimension produced by elaborate, demanding growth methods, e.g. Czochralski or Bridgman techniques [91, 98, 163, 164]. These time-consuming and expensive methods for produc ing scintillators thus impose constraints on the final radiation detection systems. Recently ceramic BGO scintill ators were reportedly produced by different methods such as solid state reaction, hy drothermal synthesis and pulse laser deposition (PLD) [163, 165-167]. The advantages of using ceramic processing for scintillators are reportedly th e homogeneous distribution of dopants, the cost-effective mass

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99 production and the versatility of sh apes and sizes [92]. In this work, we have studied flowerand coral-like BGO nanocrystals and Bi2O3/BGO core/shell scintillator composite crystals synthesized via a facile aqueous solution precipitation method. 6.2 Experimental 6.2.1 Materials Flower-like BGO crystals were synthesized via an aqueous solution precipitation process, in which 1.94 g of bismuth nitrate and 0.314 g of germanium oxide were di ssolved into 10 mL of nitric acid, followed by the addition of 9 g of ur ea with 100 mL of deionized (DI) water. The precursor was stirred and heated around 90 to have a homogeneous solution, followed by pouring of 30 mL of sodium hydroxide. After the additions of base reagent, white suspension was observed in the precursor solution. The suspension were sampling in certain time interval to study the time-dependent growth analysis. Synthe sis of coral-like BGO cr ystals were followed the same procedure except the addition of ammonia. Moreover, Core/shell Bi2O3/BGO composites were obtained by drop-wise of 30 mL of ammonia using a micro-pump in one hour and the solution stayed the same condition for an other 3 h. White precipitation was collected by washing and centrifuging with DI water several times followed by oven-dried at 60 in air. The calcination was performed at 500 in air for 2 h. 6.2.2 Characterization As-prepared and calcined samples were characterized by X-ray diffraction (XRD) (Philips APD 3720) with Cu K radiation (=1.5418 ). The XRD pattern was collected from dried powder samples in the step scan (0.02) mode over a 2 scan range of 20-70. The morphology and size of the core/shell composite were determined with field emission scanning electron microscopy (FE-SEM). Photoluminescence (P L) and photoluminescence excitation (PLE)

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100 spectra were measured at room temp erature using a JASC O FP-6500/6600 fluorescence spectrometer with a 150 W xenon lamp. The sc intillation response of the composites was characterized by means of differential pulse he ight distribution measur ements using a Hidex Triathler scintillation counter with a Ha mamatsu R850 photomultiplier (PMT) and a 241Am source (1 Ci). Samples were contained in glass veils, with background corrections based upon an empty veil. 6.3 Results and Discussion 6.3.1 Flower-Like Self-Activated BGO Crystals 6.3.1.1 Time-dependent growth analysis To investigate the morphology evolution of th e flower-like BGO, th e time dependence of growth analysis was determined by XRD and SEM. The morphology and size of as-prepared flower-like BGO colloidal crystals are shown in Figure 6-1. Based on the SEM micrographs, BGO precipitated immediately after the addition of the reagent and subsequently grew larger. As shown in Figure 6-1(a), nano-size BGO were formed after 2 min reaction time. For a reaction time of 35 min (Figure 6-1(b)), agglomerates of spherical BGO nanoparticles were observed and the growth of agglomerates continued. Subse quently, after 1 h reacti on time, colloidal BGO nanocrystals grew into micron-size pa rticles as shown in Figure 6-1(c). Figure 6-1. SEM micrographs showing the morphol ogy and particle size of as-prepared flowerlike BGO crystals at reaction times of (a) 2 min (b) 35 min and (c) 1 h.

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101 As shown in Figure 6-2(a), micron-size BGO crystals were obtained after a reaction of 1 h. Figure 6-2(b) shows a high ma gnification SEM micrograph of individual flower-like BGO crystals. The size of flower-l ike BGO structures is ~10 m and they were composed of ~1 m agglomerated crystals with sharp edges. After being calcined at 500 for 2 h in air, flower-like BGO crystal edge became rounded and a small amount of second phase was formed in the grain boundaries. This second phase is believed to be Bi12GeO20, which has a lower melting point than Bi4Ge3O12. Formation of a second phase at the grai n boundaries has been reported during laser and furnace sintering of Bi4Ge3O12 [166]. Figure 6-2. SEM micrographs of fl ower-like BGO crystals after grow th for 1 h: as-prepared at (a) low and (b) high magnification, and (c) after calcining at 500 for 2 h in air. The XRD spectra at different growth times fr om flower-like BGO crystals are shown in Figure 6-3. The spectrum in Figur e 6-3(a) indicates that as-prepared BGO nanoparticles formed in 2 min show a broad amorphous scattering peak due to insufficient time for crystallization. For a reaction of 35 min (Figure 6-3(b)), BGO crystal exhibited better crystallinity with sharper peaks and continued to improve its crystallinity with the increas ed reaction time. Sharper peaks were observed after growth for 1 h, indicating a larger crystal size (Figure 6-3(c)) of cubic Bi4Ge3O12 (JCPDS card #89-1419; Figure 6-3(d), blue pattern) without any evidence of a secondary phase.

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102 Figure 6-3. X-ray diffraction (XRD ) pattern from as-prepared flower-like BGO crystals as a function of reaction time: (a) 2 min, (b) 35 min and (c) 1h. (d) Diffraction peaks from JCPDS for cubic Bi4Ge3O12 (blue pattern). The XRD patterns from samples with diffe rent reaction times and calcined at 500 for 2 h in air are shown in Figure 6-4. Calcining resulted in observation of secondary phase of Bi12GeO20 in addition to the Bi4Ge3O12 phase, as shown in Figure 6-4(a) for growth time of 2 min. For a longer reaction time of 35 min (Figur e 6-4(b)) or 1 h (Figure 6-4(c)), the amount of Bi12GeO20 was small, which is consistent with th e small amount of secondary phase at grain boundaries in SEM micrographs for a 1 h sample (Figure 6-4(c)) and with a literature report [166].

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103 Figure 6-4. X-ray diffraction (XRD) pattern from flower-like BGO crystals calcined at 500 for 2 h in air for reaction times of (a) 2 mi n, (b) 35 min and (c) 1h. (d) Diffraction peaks from JCPDS for cubic Bi4Ge3O12 (blue pattern) and Bi12GeO20 (red pattern). 6.3.1.2 Photoluminescence analysis Figure 6-5 shows the photoluminescence (PL) and photoluminescen ce excitation (PLE) spectra of as-prepared and calcined flower-like BGO samples. A broad excitation band is distributed in the 225-300 nm range with broa d peaks at 250 and 280 nm, as shown in Figure 65(a). Trivalent bismuth has a mercury-like 6S2 electronic configuration leading to 1S03P1 and 1S01P1 transitions, which are assigned to the p eaks at ~280 and ~250 nm, respectively [161]. The emission spectra from as-prepared or calcined (400 and 600 ) flower-like BGO, and from a BGO single crystal are compared in Figure 65(b). BGO is an intrinsic self-activated scintillating material with Bi3+ ions acting as luminescent cente rs. Both flower-like and single

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104 crystal BGO exhibited a broad emission over the range from 350 to 700 nm with a maximum at 467 nm. This PL spectrum is associated with the 3P11S0 transition of Bi3+ ions. The quantum yield (QY) of scintillatio n nanoparticles is defined to be the fr actional or percentage ratio of the number of emitted photons to the number of absorbed photons [133, 168]. QYs of the three flower-like BGO samples are re ported in Figure 6-5(b). For excitation at 280 nm, the QY was ~3%, 11% and 28% from as-prepared or 400 and 600 calcined samples, respectively. Also shown in Figure 6-5(b) are color photographs of the emission exci ted by the ultra-violet (UV) spectrum from a mercury discharg e lamp. Consistent with the broad emission peak, a bluishwhite color is observed with the intensity proportiona l to the corresponding QYs. Figure 6-5. Photoluminescence spectra from flow er-like and single BGO crystal samples. (a) Photoluminescence excitation (PLE) spectru m for emission at 467 nm from calcined sample (400 2 h). (b) Photoluminescence (PL) sp ectra of as-prepared and calcined BGO polycrystalline or single BGO crystal samples excited at 280 nm. The insets show color photographs of the emissi on excited by UV irradiation and the corresponding quantum yields.

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105 6.3.2 Coral-Like Self-Activated BGO Crystals 6.3.2.1 Time-dependent growth analysis The evolution of morphology and structural phases of coral-like BGO with time of growth was analyzed by XRD and SEM. SEM microgr aphs in Figure 6-6 show the morphology evolution as a function of reacti on time from 3 min to 2 h. BGO nucleated immediately after the addition of the reagent and subse quently grew larger with time. BGO nanoparticles ~30 nm in diameter were formed in a reaction time of 2 min (Figure 6-6(a)). For a reaction time of 30 min (Figure 6-6(b)), coral-like BGO morphology was observed due to growth of agglomerated nanoparticles to produce few microns crystals. After growth for 2 h, the coral-like BGO morphology was similar to but slightly larger than the 30 min sample, as shown in Figure 6-6(c). Figure 6-6. SEM micrographs of as-prepared coral-like BGO crystals at reaction times at (a) 3 min (b) 30 min and (c) 2 h. A high magnification SEM micrograph of as-prepared coral-like BGO (Figure 6-7(a)) shows that a coral-like BGO cluster results from self-assembly of several triangular BGO nanocrystals with sizes of 200~400 nm as shown in the inset of Figure 67(a). These triangular BGO structures agglomerate to form the petal-lik e feature of the cluster. The TEM micrograph in Figure 6-7(b) shows this unique cora l-like morphology. Upon calcining at 600 for 2 h in air, the triangular BGO nanocrystals fused and the sharp edge features were more rounded, as shown in Figure 6-7(c) and (d).

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106 Figure 6-7. (a) SEM micrographs of as-prepared coral-like BGO samples after 2 h of growth. The inset shows individual nanocrystal s at a higher magnification. (b) TEM micrograph of as-prepared coral-like BGO cr ystals after 2 h growth. (c) Low and (d) high magnification SEM micrographs of coral-like BGO samples calcined at 600 for 2 h. Figure 6-8 shows XRD spectra after growth times of 3 min to 2 h, where the black spectra are from as-prepared BGO and the red spectra indi cate samples were calcined at 600 for 2 h in air. Figure 6-8(a) shows that as-prepared BGO nanoparticles formed after 3 min have a pattern corresponding to cubic Bi4Ge3O12 (JCPDS card #89-1419; Figure 6-8(d), blue pattern). After calcination, a second phase of Bi12GeO20 was observed from a 3 min sample. This minor secondary phase could result from an incomplete phase transformation du e to the short reaction time. With a reaction time of 30 min, the crys tallinity of the BGO structures improved and no secondary phase was found after the calcination (Figure 6-8(b)). For BGO crystals grown for 2 h, both as-prepared and calcined samples shows be tter crystallinity and no secondary phase was

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107 detected. These data show that crystalline colloidal BGO can be grown in aqueous solution at low reaction temperatures in short times without further heat treatment. Figure 6-8. X-ray diffraction (XRD) patterns from as-prepared (black line) and calcined (red line) coral-like BGO structures for reaction times of (a) 3 min, (b) 30 min and (c) 2h. (d) Diffraction peaks from JCPDS for cubic Bi4Ge3O12 (blue pattern) and Bi12GeO20 (orange pattern). 6.3.2.2 Photoluminescence analysis The photoluminescence (PL) and photolumines cence excitation (PLE) spectra from asprepared coral-like BGO and BGO:Eu3+ crystals as well as their CIE coordinates are shown in Figure 6-9. The excitation band for as-prepared co ral-like BGO has significant intensity between 225-300 nm with the broad peaks at ~250 and ~280 nm, as shown in Figure 6-9(a). These characteristics are similar to those for flowe r-like BGO discussed above. Assignments of the excitation peaks to transitions on the Bi3+ and the explanation for self-activated luminescence have been discussed above. These trends and expl anations are valid for both coral like as well as

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108 flowerlike BGO, and therefore will not be rep eated. Figure 6-9(b) shows the CIE coordinates (x=0.26 and y=0.32) for as-prepared coral-like BGO crystals excited at 280 nm, which matches the greenish white emission color. The phot oluminescence spectra from as-prepared Eu3+ doped BGO (BGO:Eu3+) are shown in Figure 6-9(c). The ex citation bands for doped and undoped BGO were essentially identical with the strongest peak at ~280 nm. Th e PL emission spectrum exhibit features from both the host B GO broad band emission and the Eu3+ characteristic 5D0-7Fj (J=1~4) transitions resulting in a peak at 609 nm (J=2). By doping BGO with Eu3+, the additional emission at 609 nm complements the greenish wh ite to make pure white emission with CIE coordinate of (x=0.33, y=0.33) Therefore, coral-like BGO:Eu3+ crystals have good potential as a phosphor for fluorescent lamps due to its bala nce white light emission under UV excitation. Figure 6-9. Photoluminescence (PL) and photoluminescence excita tion (PLE) spectra from asprepared coral-like (a) BGO and (c) BGO:Eu3+ crystals where PLE spectra used emission at 467 nm, and PL spectra were excited at 280 nm. The CIE coordinates from as-prepared coral-like (b) BGO and (d) BGO:Eu3+ crystals excited at 280 nm.

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109 6.3.2.3 Radiation test The scintillation properties of BGO under -ray irradiation was test ed using an americium (241Am, 1 Ci) and a cesium source (137Cs, 17 Ci). Figure 6-10 shows the differential pulse height spectra from both as-prepared and calcine d coral-like BGO crysta ls. A distinctive peak from the 59 KeV gamma from 241AM is presented in Figure 6-10(a) from as-prepared and calcined coral-like BGO samples at channel number 50 and 70, respectively. The calcination process enhanced the scintillati on brightness by ~30%, as indicat ed by increased channel number for the peak. The differential pulse height spectra in Figure 6-10(b) are the scintillation response from calcined BGO samples irradiated by 662 KeV gamma rays from the 137Cs source. For control, an empty veil was used to measur e the background. Compar ed with the background from an empty veil, the calcined coral-like BGO sample showed higher counts over a broad range. The broad range is probably a result from a large amount of scattering events in the powder sample. A number of parameters (e.g. reaction time and temperature, calcination atmosphere, temperature and time) can be furthe r tested to improve th e scintillation response. These preliminary radioluminescence data sugges t that coral-like BGO crystals synthesized by this facile aqueous pr ecipitation method show some promis e for application in radiation detection.

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110 Figure 6-10. Differential pulse height spectrum of coral-like BGO cr ystals excited by (a) 241Am and (b) 137Cs gamma-ray irradiation. 6.3.3 Bi2O3/BGO Core/Shell Composites 6.3.3.1 Time-dependent growth analysis The morphology evolution of the Bi2O3/BGO core/shell composite was determined by a time-dependent growth analysis with XRD a nd SEM. The morphology and size of as-prepared Bi2O3/BGO core/shell composite are shown in Figure 6-11. Based on SEM micrographs, rod-like Bi2O3 precipitated immediately after the addition of the reagent, while BGO nanocrystals precipitated heterogeneously as islands (Figure 611(a)) and grew laterall y to cover the surface, After 1 h, nanoagglomerates of BGO nanoparticles were observed and the size and density of the nanoagglomerates continued to increase as show n in Figure 6-11(b) and (c). With continued reaction time, BGO crystals evol ved from nanoagglomerates to a continuous dendrite-like layer on the Bi2O3 rods, forming Bi2O3/BGO core/shell composites as shown in Figure 6-11(d) and (e).

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111 Figure 6-11. SEM micrographs of Bi2O3/BGO core/shell composite at different reaction times: (a) 30 min, (b) 1 h, (c) 2 h, (d) 3 h, (e) 4 h, and (f) high magnification SEM micrograph of surface dendrite-like featur e for a 4 h reacted sample. Figure 6-12(a) shows a region of the com posite after 3 h reaction time where the Bi2O3 rods is visible and covered by a BGO layer (i dentified by energy dispersive spectrometer-EDS) and the BGO layer growing on the Bi2O3 rods is 200-300 nm thick. Bi2O3 rod was fully covered by the dendrite-like BGO shell after 4 h reaction, as shown in Figure 6-12(b). Figure 6-12. SEM micrographs of Bi2O3/BGO core/shell composite at reaction times of (a) 3 h (shell coverage incomplete), and (b ) complete shell coverage at 4 h. The XRD spectra of Bi2O3/BGO core/shell composite for reaction times from 30 min to 4 h are shown in Figure 6-13. The XRD spectrum in Figure 6-13(a) shows that the as-prepared

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112 Bi2O3 micro-rods formed after 30 min ar e poorly crystallized tetragonal Bi2O3 phase (JCPDS card #78-1793; Figure 6-13(d), or ange pattern). The XRD spect rum for as-prepared samples grown for 4 h, shown in Figure 6-13(b), is from Bi4Ge3O12 in the cubic phase (JCPDS card #891419; Figure 6-13(d), blue pattern) and the crysta llinity increases with increased reaction time. After a reaction time of 4 h plus being calcined at 500 for 2 h in air, the XRD spectrum shows the presence of both Bi2O3 and BGO as shown in Figure 6-13(c). Figure 6-13. X-ray diffraction (XRD ) spectra from as-prepared Bi2O3/BGO core/shell composite as a function of reaction time: (a) 30 min, (b) 4 h, (c) 4 h sample calcined at 500 2 h in air, and (d) diffraction peaks from JCPDS files for cubic (blue pattern) and tetragonal (orange pattern) Bi2O3. 6.3.3.2 Formation mechanism of dendrite-like shell The time dependence of morphology and structure of the Bi2O3/BGO core/shell composite suggests possible formation mechanisms. The forma tion of the core/shell composite starts with the nucleation and growth of Bi2O3 cores, followed by heterogeneous nucleation of BGO. While

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113 nucleation is generally contro lled by thermodynamic state parameters, growth is generally governed by the kinetics of processes controlled by various activation energies. These kinetic processes are formally activated and the proba bility of the process leading to growth is proportional to exp(-Ga/kT), where k is Boltzmanns constant, T is temperature and Ga is defined as the activation free ener gy [169]. Activation energies reported in th e literature for reactions that lead to the listed compounds such as Bi2O3, Bi4Ge3O12 (BGO) and GeO2 are shown in Table 6-1 [170-172]. Note that the dominant growing phase progresses from Bi2O3 to BGO in ascending order of activation free energy. In fact the phase with the highest activation energy, GeO2, is not observed. Thus it appears that the activation energy determines the progression of phases in the core/shell composite. Once the nucleation of Bi2O3 has occurred, growth quickly produces micro-rods in the aqueous colloidal precursor solution. BGO nanocrystal heterogeneously nucleated and grow on the surface of Bi2O3 cores. The dendrite-like structure of BGO nanocrystal was finally formed because the subse quent particles diffused and stuck with primary nucleated seeds. The aggregated cluster generate d by this process are highly branched which is attributed to the proliferation of instabilities induced by the diffusion-lim ited kinetic aggregation process [173-175]. Each Growth o ccurs at the dendrite-like clus ter boundaries unti l the clusters coalesce resulting in the dendrite-like shell. Table 6-1. Activation energies repo rted in the literature for select ed compounds of interest in our one-pot synthesis [173-175]. Materials Bi2O3 Bi4Ge3O12 GeO2 Activation Energy (eV) 0.55 0.9 2.3 6.3.3.3 Photoluminescent analysis The photoluminescence (PL) and photoluminescence excitation (PLE) spectra of calcined Bi2O3/BGO core/shell composite are shown in Figure 6-14. Excitation of 530 nm emission from

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114 the composite occurred over a range between 225-300 nm with the peak located at 280 nm, as shown in Figure 6-14(a). Trivalen t bismuth has a mercury-like 6S2 electronic configuration leading to 1S03P1 and 1S01P1 transitions, resulting in peaks at ~280 and ~250 nm, respectively [161], as shown in Figure 6-14(b). BGO is an intrin sic self-activated scintillator with Bi3+ ions as luminescent centers. The insets of Figure 6-14(b) shows the emission of core/shell composite under the r oom light (top) and excited by an ultra-violet (UV) mercury discharge lamp (bottom). Figure 6-14. (a) Photoluminescence excitation (emi ssion at 530 nm), and (b) emission (excitation at 280 nm) spectra from calcined Bi2O3/BGO core/shell composite. The inset shows photographs of emission irradiated by room light (top) and UV (bottom). 6.3.3.4 Radiation test The scintillation properties of the BGO under -ray irradiation was determined using an americium point source (241Am, 1 Ci). The differential pulse height spectrum of calcined Bi2O3/BGO core/shell composite is shown in Figure 6-15. A peak from the 59 KeV emission from 241Am was observed [176]. Calcined samples were better scin tillators than as-prepared samples due to better BGO crystallinity and Bi2O3 absorption. The thickness of the luminescent BGO shell of the Bi2O3/BGO core/shell composite is only a few hundred nanometers compared

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115 with a few tens of micrometers for the Bi2O3 cores. The range of a 59 KeV -ray in Bi2O3 and BGO is 5 and 4 m, respectively. Therefore, the energy cascades largely o ccur within the Bi2O3. An important property in -ray radiation detection is detector efficiency which is proportional to the -ray stopping power. The stopping power figure of merit depends on the density, effective atomic number and the thickness of the detector materials [96, 177, 178]. The Bi2O3/BGO core/shell composite possesses high stopping power and therefore is a promising scintillator. Figure 6-15. Differential pulse height spectrum of calcined Bi2O3/BGO core/shell composite irradiated by an 241Am source showing a broad scin tillation response centered at approximately channel 30. 6.4 Conclusions In our study, bismuth germanate Bi4Ge3O12 (BGO) crystals was successfully prepared by the first time through a facile precipitation met hod in aqueous solution at low temperature (~90 ) in short reaction time (~30 min). By mani pulating the experiment al conditions, various morphologies and shapes of BGO nanocrystals can be obtained, such as flower-like, coral-like and core/shell structure. Time-dependent growth analysis indicates th at morphology and size of

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116 BGO crystals evolved from ~70 nm nanoparticles af ter 3 min to coralor flower-like clusters. Well-crystallized BGO crystals were obtained after 30 min reac tion time without any post-grown heat treatment. Photoluminescent spectra of as -prepared and calcined coral-like BGO crystals showed a broad emission band with the p eak at 467 nm associated to the Bi3+ ions 3P11S0 transition similar to BGO single cr ystal. In addition, a unique Bi2O3/BGO core/shell scintillator composite was synthesized using a one-pot aqueous solution precipitation method. Timedependent analysis of the evolution of mo rphology and crystal structure showed that Bi2O3 precipitated first grew into micro-rods, follo wed by heterogeneous nucleation of BGO on the Bi2O3. The growth first of Bi2O3 followed by growth of BGO can be attributed to the relative magnitude of activation energies for processes controlling the growth of core and shell materials. The final morphology was micro-rods Bi2O3 covered with dendrite-like BGO. Photoluminescence excited by UV radiation and the core/shell scintillator response to a 241Am 59 keV -source were reported to demonstrate the promise of this ceramic scintillator for radiation detection.

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117 CHAPTER 7 ENHANCED PHOTOLUMINESCENCE OF COLLOIDAL SELF-ASSEMBLED ALMONDSHAPED GADOLINIUM VANADIUM OXIDE NANOCRYSTALS BY CO-DOPING EUROPIUM AND BISMUTH 7.1 Introduction When exposed to various excitation sources such as photons, electrons or an electric field, luminescent nanocrystals convert a fraction of the absorbed energy into characteri stic visible, ultraviolet (UV) and/or near infrared (NIR) photons [86, 179]. These luminescent materials are used in solid state lighting, X-ray detection, display and bio-labeling [59, 87, 180, 181]. In general, luminescent materials are composed of a host with a sma ll amount of dopant ions, called activators. Host matrices are usually oxides, su lfides, oxysulfides or hali des. Rare-earth (RE) dopant ions (activators) are of ten used to control the charact eristic emission wavelength and intensity [125, 182]. RE doped orthovanadate phos phors, widely applied in lamps and displays, have been shown to have high luminescent e fficiency under UV excitation [183, 184]. Co-doped Bi3+ has been used as sensitizer ions to capture and transfer energy to Eu3+ activators in several host matrices [78, 185-187]. Bulk YVO4 or GdVO4 are generally synthesized with complex solid-state reactions that require times of seve ral hours to several days at temperatures above 1000 [72, 186, 188, 189]. These syntheses proce dures often result in inhomogeneous doped micrometer-size agglomerates that contain seco nd impurity phases which lead to luminescent intensities too low for biological and high-definiti on flat display applications. As a result, other synthesis methods, such as combustion, hydrothe rmal and sol-gel methods have been reported [81, 185, 190, 191]. However, the synthesis co nditions reported are non-hydrolytic solutions with reaction temperatures higher than 170 for several hours, or aqueous solution resulting in micrometer-sized crystals. In this study, colloidal well-crystallized GdVO4 nanocrystals codoped with Eu3+ and Bi3+ were prepared in a water-based solution without surfactants at low

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118 reaction temperatures (< 90 ) in short times (3 min). In addition, dried GdVO4 powder can be re-dispersed in water forming clear colloidal solution. The structural and photoluminescent properties of GdVO4 nanocrystals co-doped with Eu3+ and Bi3+ were also studied. 7.2 Experimental 7.2.1 Materials All of the reagents and chemicals were used as received without further preparation. For synthesis of GdVO4 nanocrystals, a water-based solution precipitation method was used without any surfactant. Specifically, 0.03 mol/L of Gd(NO3)3 and 0.003 mol/L of Eu(NO3)3 were dissolved into 100 mL of deionized (DI) wa ter, followed by addition of 5 mL of HNO3. For most samples, various amounts of Bi(NO3)3 was also added to the precursor solution to achieve doping concentrations between 2~ 15 % with respect to the Gd3+ concentration. The clear aqueous solutions were subs equently heated at 85~90 and 0.03 mol/L of NH4VO3 was added under vigorous stirring. After th e additions of 30 mL of NH4OH, the precursor solution changed from a transparent yellowish color to a vivid yell ow, and finally to a translucent white color for reaction times of 3 min to 2 h. The white preci pitate was collected by washing and centrifuging with DI water several times. Powder of lu minescent nanocrystals were oven-dried at 60 in air, but could be re-dispersed making a clear colloi dal aqueous solution. As-prepared colloidal and dried powder GdVO4:Eu3+, Bi3+ nanocrystals showed intense photoluminescent (PL) emission under ultra-violet (UV) excitati on from a mercury discharge lamp. 7.2.2 Characterization The crystal structure of as-prepared nanoc rystals was determined by X-ray diffraction (XRD) (Philips APD 3720) with Cu K radiation source (=1.5418 ). The XRD pattern was collected from dried powder samples in the step scan (0.02) mode over a 2 scan range of 20-70.

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119 The morphology and size of the nanoparticles were determined by field emission scanning electron microscopy (FE-SEM) and high reso lution transmission electron microscopy (HRTEM). Photoluminescence (PL) and photoluminescence excitation (PLE) spectra for the colloidal solutions were measured at room temperature using a JASCO FP-6500/6600 fluorescence spectrometer with a 150 W xenon lamp. 7.3 Results and Discussion 7.3.1 Structure and Morphology of GdVO4:Eu3+/Bi3+ Nanocrystals Figure 7-1 shows the XRD pattern from dried as-prepared GdVO4:Eu3+ nanocrystals for a reaction time of 3 min. Figure 7-1(a) shows that as-prepared nanoc rystals were crystalline. An average crystal diameter of 10 nm was calculated from Scherrers equation from a 3 min sample, which is consistent with TEM data reported below (Figure 7-3). After being calcined at 800 for 2 h in air, the XRD peaks became more intense and sharper, indicating larger, crystalline nanocrystals. The diffraction patterns from as-p repared and calcined samples correspond to the cubic phase GdVO4 (JCPDS card #81-0563 shown in Figure 7-1(c)). SEM micrographs of asprepared nanocrystals after reaction for 3 min and 2 h are shown in Figure 7-2(a) and (b), respectively. The dried nanocrystals tend to a gglomerate with a size of 50-70 nm. After being calcined at 800 for 2 h, the GdVO4 agglomerates grew to micrometer size crystals, as shown in Figure 7-2(c) and (d).

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120 Figure 7-1. X-ray diffraction (XRD ) patterns from (a) as-prepared, (reacted for 3 min. at <90 ), (b) calcined (800 2 h in air) GdVO4:Eu3+ nanocrystals, and (c) diffraction peaks from JCPDS (#86-0996) for cubic GdVO4. Figure 7-2. SEM micrographs of as-prepared GdVO4:Eu3+ nanocrystals after reaction times of (a) 3 min and (b) 2 h, and calcined samples (800 2 h in air) reacted for (c) 3 min and (d) 2 h prior to calcining.

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121 Figure 7-3(a) and (d) show low to high magnification TEM micrographs,of as-prepared GdVO4 nanocrystals grown for 3 min. Non-agglomerated individual GdVO4 nanorod ~10 nm in diameter (Figure 7-3(c)) self assemble into al mond-shaped clusters ~50 nm wide and ~100 nm long (Figure 7-3(a) and (b)). High-resolution TEM of the GdVO4 nanocrystals in Figure 7-3(d) shows well-resolved crystalline fringes, consistent with the spotty diffraction rings in the inset of selected area electron diffraction pattern (SAED). The spacing of ~3.6 between lattice fringes corresponds to the interplanar spacing of the (200) planes of cubic GdVO4. Figure 7-4(a)-(d) show TEM micrographs of as-prepared GdVO4 nanocrystals grown for 2 h. Almond-shaped nanoclusters composed of self-assembled nanor ods are again observed. The size and morphology of both the almond-shaped clusters and nanorods are constant with reaction times up to 2 h. Figure 7-3. TEM micrographs of GdVO4:Eu3+ nanocrystals reacted for 3 min showing almondshaped clusters self-assembled from nanorods. Note the change from low to high magnification in (a)-(c). (d) High-reso lution TEM image showing well-resolved lattice fringes. Inset: selected area electron diffraction pattern (SAED).

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122 Figure 7-4. TEM micrographs, at progressive ly higher magnification in (a)-(d), of GdVO4:Eu3+ nanocrystals after 2 h of growth. Note that the self-assembled almond-shaped clusters were of the same size (30~60 nm) as those grown for 3 min. 7.3.2 Photoluminescent Properties of GdVO4:Eu3+/Bi3+ Nanocrystals Room temperature photoluminescence excitation (PLE) spectra from as-prepared colloidal almond-shaped GdVO4 nanocrystals grown for 2 h and co-doped with 10 mol% Eu3+ and 0-15 mol% Bi3+ are shown in Figure 7-5(a). For GdVO4:Eu3+ nanocrystals, the emission at 617 nm is assigned to the 5D0-7F2 transition of Eu3+. The weak, sharp PLE peak at 396 nm is from the selfactivated 5F0-5L6 Eu3+ transition. The broad excitation p eak from 225 to 375 nm is from two charge transfer bands [187]. The two charge transfer bands resu lt from O-Eu and VO4-Eu charge transfer processes centered at ~280 and ~320 nm, respectively [ 61, 81]. The summed excitation peak from the charge transfer ba nds has a peak at ~323 nm with the band edge at ~360 nm. With the Eu3+ concentration fixed at 10%, increasing the Bi3+ concentration from 2% to 15% shifted the charge transfer band edge to longer waveleng th (by as much as 30 nm). Color photographs in

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123 Figure 7-5(b) shows the color and relative inte nsity of photoluminescence from as-prepared GdVO4:Eu3+(10%), and GdVO4:Eu3+(10%), Bi3+(2%) nanocrystals excited by a mercurydischarge lamp with a spectral distribution labe led either as long or short wavelength. Under short and long UV wavelength excita tion, both samples show high intensities, but because of the extension of the excitati on to longer wavelength by Bi3+ ions, co-doped samples show perceptibly brighter emissi on than that from only Eu3+ doped samples. This extension in the excitation band is attributed to additional abso rption from the Bi-O bonds and the subsequent energy transfer between O and Eu [185, 192, 193]. As shown in Figure 7-5(a), increased concentrations of Bi3+ led to an increase in excitation from the O-Eu band located at ~280 nm relative to the V-Eu charge tran sfer band located at ~320 nm. Ther efore, the predominant charge transfer band of GdVO4:Eu3+ is the V-Eu band, while the O-Eu band is dominant for concentrations of Bi3+ > 7%. Figure 7-5. (a) Room temperature photolumines cence excitation (PLE) spectra from as-prepared GdVO4:Eu3+ (10%) nanocrystals grown for 2 h, and with various concentrations of co-doped Bi3+ ( emission = 617 nm). (b) Colored photogra phs showing the intensity and color of emission from GdVO4:Eu3+ and co-doped GdVO4:Eu3+, Bi3+ nanocrystals excited by a handheld mercury discharge lamp that emitted ultraviolet (UV) photons over wavelength regions denoted as a short or long spectrum [194].

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124 Photoluminescence (PL) of as-prepa red 2h growth almond-shaped GdVO4 nanocrystals codoped with Eu3+ and Bi3+ are shown in Figure 7-6(a)-(c) for excitation at 280, 323 and 347 nm, respectively. The inset in Figure 7-6(a) s hows the bright red em ission from colloidal GdVO4:Eu3+ nanocrystals in a quartz cuvette unde r ~280 nm UV excitation. Generally, the photoluminescence spectrum of Eu3+ doped crystals exhibits str ong, sharp peaks at 595 and 617 nm due to the 5D0-7F1 and 5D0-7F2 transitions, respectively [57]. The spectra for three different excitation wavelengths show identical peak wavelengths from Eu3+. However, the PL intensity from GdVO4:Eu3+, Bi3+ nanocrystals excited by 280 and 323 nm, shown in Figure 7-6(a) and (b), respectively, were enhan ced by 45-90% for 2l% of Bi3+ versus singly doped GdVO4:Eu3+(10%) nanocrystals. The PL intensity decreased for concentrations of Bi3+ >10%. The changes in the 5D0-7F2 transition peak intensity wi th the concentration of Bi3+ are consistent with the changes in the excitation spectra shown in Figure 7-5(a). The change in PL intensity from co-doped nanocrystals is more significant when excited by longer wavelengths, e.g. 347 nm as shown in Figure 7-6(c), due to the extens ion of excitation band at longer wavelengths. The PL from nanocrystals was improved by ~5.7 times wh en excited 347 nm by incorporating 2% Bi3+ ions. The PLE and PL data show that the sensitization effect of Bi3+ on Eu3+ emission is dependent on the concentration of Bi3+, with maximum PL intensity for 2% and decreased PL intensity for higher concentrations. The decreased PL intensity for Bi >7 mol% presumably results from increased non-radiative relaxations due to c oncentration quenching or formation of trapping centers due to a high concentrations of aggregated Bi3+ ions [189, 195].

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125 Figure 7-6. Photoluminescence (PL) spectra of as-prepared GdVO4:Eu3+ (10%) nanocrystals (grown for 2 h) co-doped with concentrations of Bi3+ ranging from 0-15%, with an excitation wavelength of (a) 280 nm (Inse t: emission from nanocrystals in a quartz cuvette), (b) 323 nm and (c) 347 nm. 7.4 Conclusions Highly crystallized GdVO4:Eu3+, Bi3+ nanocrystals were prepared by a facile water-based solution precipitation met hod at temperatures <90 in a time short as three minutes and without surfactants or additional solvents .. Individual nanocrysta ls self assembled into nanorods which agglomerated into almond-shaped GdVO4 clusters with dimension of 10 nm wide by 100 nm long. Intense PL emission can be observed from as prepared GdVO4:Eu3+, Bi3+ nanocrystals as

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126 either solution or dried powder. Furthermore the PL intensity from GdVO4:Eu3+ nanocrystals codoped with 2% Bi3+ was increased by factors rang ing from 45% to 570%, depending upon excitation wavelength. The sensitization of Eu by Bi resulted from the additional absorption of Bi-O component extending the excitati on band to longer wavelengths. With Bi3+ concentrations >10%, the PL intensity decreased presumably due to concentration quenching and charge trapping.

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127 CHAPTER 8 CONCLUSIONS AND FUTURE WORK 8.1 Luminescent Zero and One Dimensional Gd2O3:Eu3+ Nanocrystals In this study, well disperse d colloidal crystalline Gd2O3:Eu3+ nanocrystals were synthesized at 290 by a hot solution non-hydrolytic method. In our study, two distinct shapes of nanocrystals were prepared, e. g. flower-like an d sphere-like Gd2O3:Eu3+ nanocrystals, by manipulating the reaction temper ature and heating se quence. Photoluminescence excitation (PLE) spectrum for emission at 609 nm from these as-prepared Gd2O3:Eu3+ nanocrystals showed a broad excitation band from 230 to 300 nm which is the charge transfer band (CTB) from O2to Eu3+. The red emission from Eu3+ is dominated by the 5D0-7F2 transitions at 609 nm with minor peaks from 5D0-7Fj (j=0, 1, 3, 4) characteristic transitions. Rod-like luminescent Gd2O3:Eu3+ nanocrystals with diameters of 70 nm and lengths of 800 nm were synthesized via a facile template-free sol-gel precipitation met hod at a low reaction temperature (~85 ) followed by calcining at 800 for 2 h in air. The addition of PEG-8000 apparently led to slower growth along the [001] direction, and cons equently led to the shorter na no-rods with a smaller aspect ratio 2. 8.2 Eu3+ Doped SiO2/Gd2O3 Core/Shell Scintillation Nanoparticles Core/shell nanostructures with mono-dispersed 220 nm SiO2 cores and ~13 nm Gd2O3 shells were prepared by a solution precipitati on method. Photoluminescence (PL) was achieved by doping either the SiO2 core or Gd2O3 shell with Eu3+ ions. A thin Gd2O3:Eu3+ shell exhibited good PL with a quantum yield ( QY) up to 25 %. Moreover, SiO2/Gd2O3:Eu3+/X core/shell/shell nanostructures were synthesized with X being undoped SiO2 or Gd2O3. The amorphous SiO2 second shell decreased the QY to 16.5%, while the crystalline Gd2O3 second shell increased the QY to 32.2%. In addition, an undoped Gd2O3 shell on a SiO2:Eu3+ core increased the QY from

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128 ~2.2 % for the bare core to a maximum of ~8.5 %. The increased PL intensity was attributed to the crystalline Gd2O3 layer serving as a sensitizer with ener gy transfer to the Eu3+ either in the Gd2O3:Eu3+ shell or SiO2:Eu3+ core, along with the passivation of non-radiat ive surface defect states. 8.3 Self-Assembled Hierarchical Bi2O3/Gd2O3:Eu3+ A unique hierarchical 3D self-assembled flower-like -phase Bi2O3 was produced in the presence of PEG-8000 via a facile solution precip itation process without any templates. This is the first report of the gr owth of body centered cubic -phase Bi2O3 at such a low temperature (85 ) in short reaction time (45 min). By varying the reaction time and PEG-8000 capping agent concentration, the evolution of morphology and cr ystal phase of bismuth oxide was investigated. During the initial reaction, bismuth oxide nucleated and grew to nano-size particles in < 1 min. With continued reaction, nano-size particles were oriented and attached to other nano-particles, leading to agglomerated nano-particles. The PEG chains result in the morphology of the initially formed bismuth oxide clusters and leads to bri dging flocculation which binds triangular platelets into individual petal-like struct ures, resulting in flower-like Bi2O3. Further deposition of europium doped gadolinium oxide on this hierar chical bismuth oxide leads to a core/shell scintillator composite. The composite exhibits the predominant 5D0-7F2 transition at 609 nm upon calcination indicating a good coverage of gadolinium oxide layer and the potential application in the scintillation detector. 8.4 Self-Activated Bi4Ge3O12 Scintillator Materials Bismuth germanate Bi4Ge3O12 (BGO) crystals were successfully prepared by the first time through a facile precipitation method in a queous solution at low temperature (~90 ) in short reaction time (~30 min) with various morphologies and shapes, su ch as flower-like, coral-like and core/shell structure. Time-dependent growth analysis indicated th at the morphology and size

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129 of BGO crystals evolved from ~ 70 nm nanoparticles after 3 min to coralor flower-like clusters after 2 h. Well-crystallized BGO crystals were obtained after 30 min growth without any postgrown heat treatment. Photoluminescent spectr a of as-prepared and calcined BGO crystals showed a broad emission band with the p eak at 467 nm associated with the Bi3+ ions 3P11S0 transition, similar to a BGO single crystal. In ad dition, time-dependent anal ysis of the evolution of morphology and crystal structure of a unique Bi2O3/BGO core/shell scintillator composite showed that Bi2O3 precipitated first grew into micro-r ods, followed by heterogeneous nucleation of BGO on the Bi2O3. The growth first of Bi2O3 followed by growth of dendrite-like BGO can be attributed to the relative magnitude of activati on energies for processes controlling the growth sequence of core and shell mate rials. PL spectra excited by UV radiation and the core/shell scintillator response to a 241Am 59 keV -source were reported to demonstrate the promise of this ceramic scintillator for radiation detection. 8.5 Enhanced Photoluminescen ce of Self-Assembled Gd2O3:Eu3+, Bi3+ Nanocrystals Colloidal GdVO4 nanocrystals were synthesized by a facile water-based solution precipitation method at 90 without surfactant and other solvent. Highly-crystallized selfassembled almond-shape GdVO4 nanocrystals composed of individual nano-rods with the dimension of 10 nm wide and 100 nm long were obtained after 3 min grow th without post-grown heat treatment. The intensity of photoluminescence from GdVO4:Eu3+ nanocrystals was increased by co-doping Bi3+ with the concentration between 2-10%, while the PL intensity decreased for higher concentration of Bi3+ ions, e.g. >10%. The excitation band edge of co-doped nanocrystals extended to longer wavelengths due to the addi tional absorption from Bi-O component. With Bi3+ concentrations >10%, the PL inte nsity decreased presumably due to concentration quenching and charge trapping leading to the non-radiative relaxation.

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130 8.6 Future Work Based on the results of this research, there are a number of topics requiring work in order to demonstrate low cost, large area and mass produced ceramic scinti llators. For example, radioluminescence test should be used to ev aluate the scintillati on properties of SiO2/Gd2O3:Eu3+ core/shell (Chapter 4) and GdVO4:Eu3+, Bi3+ codoped nanocrystals (Chapter 7). The shell thickness in core/shell nanostr ucture should be optimized for surface passivation and energy transfer. Other hetero-shell cappi ng materials should be tested for the degree of enhancement of light output. For better scintillation under X-ray or -ray irradiation, the dimensions of Bi2O3/BGO core/shell structures should be matched to the attenuation distance of the X-rays and -rays so that the deposited energy needs will lead to increased lumine scence from the BGO shell through increased absorbance by the Bi2O3 cores (Chapter 6). In all ca ses, the effects of calcination temperature, atmosphere and time on the scinti llation response should be further studied. The quantum yield of scintillation nanocrystals excited by UV photons and radioactive decay particles should be measured and compared with the single crystal materials. The mechanisms leading to different scintillation intensities for excitation by UV or by radi oactive decay particles should be investigated. Finally, the major ques tion in the use of nanoc rystals for scintillation detectors remains the scattering of luminescence by the particles, which lead s to lower intensities and poor energy resolution. N ovel ideas for reducing scatte ring and for incorporating nanocrystals into a matrix are required in order to make a detector.

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BIOGRAPHICAL SKETCH Teng-Kuan Tseng was born in Kaohsiung, Taiwan in 1975. He entered the Department of Mechanical Engineering at National Central Univ ersity in 1994 and expl ored his interest in materials science. He continued his study in National Taiwan Univer sity of Science and Technology and obtained his Master of Science in 2000. He investigated magnetic thin film and characterized them using in-situ surface characterization tools in an ul tra high vacuum (UHV) system. After graduation, he joined Research Cent er for Applied Center, a nd Institute of Physics in Academic Sinica as a research associate unt il 2005. During the five periods, he immersed himself in multiple research projects, worked wi th dynamic model simulation, and fabricated and tested prototype spray devices. He construc ted from scratch a custom vacuum thermal evaporation system and synthesized various nano particles and characteri zed their structural, optical and physical properties. He was admitted to the Department of Materials Science and Engineering at the University of Florida in 2005 to pursue his Ph.D. degree. He initiated to construct the facilities for and de veloped an expertise in fabricat ion and testing of inorganic and primarily organic solar cells. His current research topics in Prof. Holloways group are synthesis and characterization of luminescen t nanocrystals for application in solar cells, solid state lighting and scintillator radiation detect ors. He received his Doctor of Philosophy in 2010 in materials science and engineering.